Fe-based metal sheet and manufacturing method thereof

ABSTRACT

A cast slab containing C: less than 0.02 mass % and made of an Fe-based metal of an α-γ transforming component is subjected to hot rolling at a temperature of an A3 point or higher and is subjected to α-region rolling at a temperature of 300° C. or higher and lower than the A3 point, and thereby a base metal sheet having a {100} texture in a surface layer portion is fabricated. Then, by performing a heat treatment under predetermined conditions, an Fe-based metal sheet is obtained in which a Z value is not less than 2.0 nor more than 200 when intensity ratios of respective {001}&lt;470&gt;, {116}&lt;6 12 1&gt;, and {223}&lt;692&gt; directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z=(A+0.97B)/0.98C is satisfied.

This application is a Continuation of application Ser. No. 14/114,131, filed on Oct. 25, 2013, which is the National Stage Entry of PCT International Application No. PCT/JP2012/061385, filed on Apr. 27, 2012, which claims priority under 35 U.S.C. §119(a) to Japanese Patent Application No. 2011-100014, filed in Japan on Apr. 27, 2011, Japanese Patent Application No. 2011-101893, filed in Japan on Apr. 28, 2011, and Japanese Patent Application No. 2012-070166, filed in Japan on Apr. 26, 2012, all of which are hereby expressly incorporated by reference into the present application.

TECHNICAL FIELD

The present invention relates to an Fe-based metal sheet having a high accumulation degree of {200} planes suitably used for magnetic cores and the like of electric motors, power generators, and transformers and capable of contributing to downsizing of these magnetic cores and reduction in energy loss, and a manufacturing method thereof.

BACKGROUND ART

Electrical steel sheets alloyed with silicon or/and the like have been conventionally used for magnetic cores of electric motors, power generators, transformers, and the like. Among electrical steel sheets, non-oriented electrical steel sheets having relatively random crystal orientations can be manufactured at a low cost, to thus be used for motors, transformers, and the like of home electric appliances, and the like in a multipurpose manner. The crystal orientations of this non-oriented electrical steel sheet are random, thus making it impossible to obtain a high magnetic flux density. In contrast to this, grain-oriented electrical steel sheets having aligned crystal orientations can obtain a high magnetic flux density, to thus be applied to high-end use for driving motors and the like of HV vehicles and the like. However, in a manufacturing method of a grain-oriented electrical steel sheet that is industrialized currently, a long-time heat treatment is required, to thus increase the cost.

As above, in the non-oriented electrical steel sheet, a sufficiently high magnetic flux density cannot be obtained, and in the grain-oriented electrical steel sheet, the direction in which a high magnetic flux density can be obtained is limited to one to two direction/directions. On the other hand, in HV vehicles, and the like, achievement of high torque and downsizing are required, and there is a demand for manufacturing a metal sheet capable of obtaining a high magnetic flux density in an in-plane circumferential direction thoroughly as a metal sheet to be used for core materials of driving motors, and the like. Thus, as methods other than the industrialized manufacturing method of the grain-oriented electrical steel sheet, there have been proposed a technique of increasing an accumulation degree of a specific crystal orientation and various techniques of decreasing a core loss. However, in the technique described in Patent Literature 7, for example, it is possible to increase an accumulation degree of {200} planes, but directionality to a specific orientation occurs, to thus have a high magnetic flux density in a specific direction, but a high magnetic flux density cannot be obtained in an in-plane circumferential direction thoroughly, and the like, resulting in that in a conventional technique, satisfactory properties are not necessarily obtained.

CITATION LIST Patent Literature

Patent Literature 1: Japanese Laid-open Patent Publication No. 10-168542

Patent Literature 2: Japanese Laid-open Patent Publication No. 2006-45613

Patent Literature 3: Japanese Laid-open Patent Publication No. 2006-144116

Patent Literature 4: Japanese Laid-open Patent Publication No. 10-180522

Patent Literature 5: Japanese Laid-open Patent Publication No. 01-252727

Patent Literature 6: Japanese Laid-open Patent Publication No. 07-173542

Patent Literature 7: International Publication Pamphlet No. WO2011/052654

SUMMARY OF INVENTION Technical Problem

Thus, an object of the present invention is to provide an Fe-based metal sheet that is likely to become magnetized in a sheet plane and further has a texture capable of obtaining a high magnetic flux density thoroughly in an in-plane circumferential direction, and a manufacturing method thereof.

Solution to Problem

The present inventors, as a result of earnest examination, found that an orientation ratio to a specific orientation is controlled with respect to an Fe-based metal of an iron sheet or the like, and thereby a <100> orientation in αFe is more densely and thoroughly distributed in a metal sheet plane to make it possible to obtain a high magnetic flux density thoroughly in an in-plane circumferential direction.

Further, the present inventors conceived that in order to manufacture such an Fe-based metal sheet, a texture in which an accumulation degree of {100} planes is increased is first formed in a surface layer portion, and at the time of γ-α transformation by the subsequent heat treatment, the texture is transformed while taking over its {100} texture. Then, they earnestly examined a method of forming the {100} texture in the surface layer portion and achievement of high accumulation of {200} planes using the γ-α transformation.

As a result, it was found that when the Fe-based metal sheet is manufactured from a slab by rolling, a rolling temperature and a reduction ratio are optimized, thereby making it possible to form the {100} texture in at least the surface layer portion. Then, it was found that when the {100} texture in the surface layer portion is taken over by using the γ-α transformation thereafter, a different metal except Fe is made to diffuse beforehand from the surface and a diffused region is turned into an α-Fe phase, and thereby in the region turned into the α-Fe phase, the {100} texture is formed, and at the time of the γ-α transformation, an accumulation degree of {200} planes in the α-Fe phase further generated by the transformation increases and the <100> orientation is distributed more densely and thoroughly, thereby making it possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly.

Further, the present inventors found that in the case of a large amount of C content being contained, when the C content is decreased by decarburization annealing, the decarburization annealing is performed under predetermined conditions, thereby also making it possible to form the {100} texture in at least the surface layer portion, and in the Fe-based metal sheet obtained finally, the <100> orientation is distributed more densely and thoroughly, thereby making it possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly.

The gist of the present invention made as a result of such examinations is as follows.

(1) An Fe-based metal sheet, includes: at least one type of ferrite-forming element except Fe, in which when intensity ratios of respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z=(A+0.97B)/0.98C is satisfied, a Z value is not less than 2.0 nor more than 200.

(2) The Fe-based metal sheet according to (1), in which the ferrite-forming element diffuses from a surface to be alloyed with Fe.

(3) The Fe-based metal sheet according to (1) or (2), in which a layer containing the ferrite-forming element is formed on at least one side of surfaces of the Fe-based metal sheet, and the ferrite-forming element that has diffused from part of the layer is alloyed with Fe.

(4) The Fe-based metal sheet according to (3), in which a thickness of the layer containing the ferrite-forming element is not less than 0.01 μm nor more than 500 μm.

(5) The Fe-based metal sheet according to any one of (1) to (4), in which an accumulation degree of {200} planes is not less than 30% nor more than 99%, and an accumulation degree of {222} planes is not less than 0.01% nor more than 30%.

(6) The Fe-based metal sheet according to any one of (1) to (5), in which the ferrite-forming element is one type of element or more selected from a group consisting of Al, Cr, Ga, Ge, Mo, Sb, Si, Sn, Ta, Ti, V, W, and Zn.

(7) The Fe-based metal sheet according to any one of (1) to (6), in which at least a partial region including the surfaces of the Fe-based metal sheet is an α single phase region made of an α single phase based component, and a ratio of the α single phase region to a cross section of the Fe-based metal sheet is 1% or more.

(8) The Fe-based metal sheet according to any one of (1) to (7), in which a thickness of the Fe-based metal sheet is not less than 10 μm nor more than 6 mm.

(9) The Fe-based metal sheet according to any one of (1) to (8), in which the α single phase region is formed on a front surface side and a rear surface side of the Fe-based metal sheet, and a crystal grain straddling the α single phase region on the front surface side and the α single phase region on the rear surface side is formed.

(10) A manufacturing method of an Fe-based metal sheet, includes:

performing hot rolling on a cast slab containing C: less than 0.02 mass % and made of an Fe-based metal of an α-γ transforming component at a temperature of an A3 point of the cast slab or higher to obtain a hot-rolled sheet;

performing α-region rolling on the hot-rolled sheet at a temperature of higher than 300° C. and lower than the A3 point of the cast slab to obtain a rolled sheet;

performing cold rolling on the rolled sheet to obtain a base metal sheet having a thickness of not less than 10 μm nor more than 6 mm;

bonding a ferrite-forming element to one surface or both surfaces of the base metal sheet;

heating the base metal sheet having had the ferrite-forming element bonded thereto up to an A3 point of the base metal sheet; and

further heating the heated base metal sheet to a temperature of not lower than the A3 point of the base metal sheet nor higher than 1300° C. and holding the base metal sheet; and

cooling the heated and held base metal sheet to a temperature of lower than the A3 point of the base metal sheet.

(11) The manufacturing method of the Fe-based metal sheet according to (10), in which a reduction ratio in the α-region rolling is −1.0 or less in terms of true strain, and the sum of the reduction ratio in the α-region rolling and a reduction ratio in the cold rolling is −2.5 or less in terms of true strain.

(12) The manufacturing method of the Fe-based metal sheet according to (10) or (11), in which

a reduction ratio in the hot rolling is −0.5 or less in terms of true strain.

(13) A manufacturing method of an Fe-based metal sheet, includes:

heating a steel sheet containing C: not less than 0.02 mass % nor more than 1.0 mass %, having a thickness of not less than 10 μm nor more than 6 mm, and made of an Fe-based metal of an α-γ transforming component to a temperature of an A1 point or higher and a temperature at which a structure is turned into an α single phase when decarburization is performed until C becomes less than 0.02 mass %, to obtain a base metal sheet that has been subjected to decarburization in a range of not less than 5 μm nor more than 50 μm in a depth direction from its surface until C becomes less than 0.02 mass %;

bonding a ferrite-forming element to one surface or both surfaces of the base metal sheet;

heating the base metal sheet having had the ferrite-forming element bonded thereto up to an A3 point of the base metal sheet; and

further heating the heated base metal sheet to a temperature of not lower than the A3 point of the base metal sheet nor higher than 1300° C. and holding the base metal sheet; and

cooling the heated and held base metal sheet to a temperature of lower than the A3 point of the base metal sheet.

(14) The manufacturing method of the Fe-based metal sheet according to (13), in which the steel sheet made of the Fe-based metal further contains Mn of 0.2 mass % to 2.0 mass %, and decarburization and demanganization are performed in a combined manner.

(15) The manufacturing method of the Fe-based metal sheet according to (13) or (14), further includes:

performing carburization on a steel sheet containing C: less than 0.02 mass %, having a sheet thickness of not less than 10 μm nor more than 6 mm, and made of an Fe-based metal of an α-γ transforming component to control C to not less than 0.02 mass % nor more than 1.0 mass %.

Advantageous Effects of Invention

According to the present invention, it is possible to manufacture an Fe-based metal sheet capable of obtaining a high magnetic flux density thoroughly in an in-plane circumferential direction.

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 is a view for explaining a method of calculating an average magnetic flux density B50;

FIG. 2 is a conceptual diagram showing the relationship between a Z value and a ratio B50/Bs of the average magnetic flux density B50 to a saturation magnetic flux density Bs and a magnetic flux density difference ΔB;

FIG. 3A is s view schematically showing a structure of a cross section of a base metal sheet having a {100} texture formed in a surface layer portion;

FIG. 3B is a view schematically showing the structure of the cross section of the base metal sheet having a different metal layer formed in the surface layer portion;

FIG. 3C is a view schematically showing the structure of the cross section of the base metal sheet in a temperature increasing process;

FIG. 3D is a view schematically showing the structure of the cross section of the base metal sheet in a heating and holding process;

FIG. 3E is a view schematically showing the structure of the cross section of the base metal sheet in a cooling process;

FIG. 4A is a view schematically showing the structure of the cross section of the base metal sheet in a state of being held at a temperature of an A3 point or higher;

FIG. 4B is a view schematically showing the structure of the cross section of the base metal sheet after cooling in the case when the different metal layers are made to remain;

FIG. 4C is a view schematically showing the structure of the cross section of the base metal sheet in the case when the base metal sheet is alloyed up to its center portion in a state of being held at the temperature of the A3 point or higher;

FIG. 4D is a view schematically showing the structure of the cross section of the base metal sheet after cooling in the case when the base metal sheet is alloyed up to the center portion; and

FIG. 5 is a view schematically showing the structure of the cross section of the base metal sheet in which a crystal grain becomes coarse.

DESCRIPTION OF EMBODIMENTS

Generally, an orientation of easy magnetization exists in α-Fe crystal, and when in a direction in which direction cosines between <100>, <010>, <001> orientations, (which will be called a [100] orientation generically), and the orientation are large, excitation is performed in a fixed magnetic field and magnetometry is performed, a high magnetic flux density is likely to be obtained. On the other hand, when in a direction in which direction cosines with respect to a <111> orientation being an orientation of hard magnetization are large, excitation is performed and magnetometry is performed, a high magnetic flux density is unlikely to be obtained. The present inventors found that more [100] orientations in the α-Fe crystal exist in a sheet plane and further the α-Fe crystal is controlled to a specific texture that is thoroughly distributed in the sheet plane, and thereby direction cosines with respect to the [100] orientation always become large in an arbitrary direction in the metal sheet plane, and when a magnetic field is applied in an arbitrary direction in the metal sheet plane and magnetometry is performed, a high magnetic flux density can be obtained.

It is characterized in that a specific texture that an Fe-based metal sheet of the present invention has contains at least one type of ferrite-forming element except Fe, in which when intensity ratios in respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z=(A+0.97B)/0.98C is satisfied, a Z value is not less than 2.0 nor more than 200.

Next, the previously described Z value will be explained.

The main orientations on which attention is focused in the present invention are {001}<470>, {116}<6 12 1>, and {223}<692>. When examining the state of a three-dimensional texture calculated by a vector method, the present inventors noticed that X-ray random intensity ratios in the above-described three plane orientations change depending on a magnetic property of a product, and learned that mathematizing this makes it possible to quantify the relationship with a magnetic property of a product and reached the present invention.

The X-ray random intensity ratios of these respective orientations may be obtained from a three-dimensional texture calculated by a vector method based on a pole figure of {110}, or may also be obtained from a three-dimensional texture calculated by a series expansion method using a plurality (preferably three or more) of pole figures out of pole figures of {110}, {100}, {211}, and {310}. For the X-ray random intensity ratios in the above-described respective crystal orientations by the latter method, for example, intensities of (001)[4 −7 0], (116) [1 −12 1], and (223) [6 −9 2] at a φ2=45° cross-section of the three-dimensional texture may be used as they are.

Subsequently, there will be explained a reason for which the expression of Z=(A+0.97B)/0.98C was found.

First, the intensity of the {001}<470> orientation is set to A. This orientation is in the {100} plane, so that direction cosines with respect to the {100} plane are 1.0. In the {100} plane, the [100] orientation being the orientation of easy magnetization exists, and thus orientation of this plane in the metal sheet plane is advantageous for obtaining a high magnetic flux density in the metal sheet plane. Thus, the intensity A is weighted with the direction cosines of 1.0 in terms of the degree of contribution to improving a magnetic flux density to be set to one of parameters in the Z value.

Next, the intensity of the {116}<6 12 1> orientation is set to B. An angular difference between this orientation and the {001} plane is 13.3° and direction cosines are 0.97. In the {001} plane as well, the [100] orientation being the orientation of easy magnetization exists, and thus orientation of this plane in the metal sheet plane is advantageous for obtaining a high magnetic flux density in the metal sheet plane. For this reason, the intensity B is weighted with the direction cosines of 0.97 in terms of the degree of contribution to improving a magnetic flux density to be set to one of parameters in the Z value.

Further, the intensity of the {223}<692> orientation is set to C. An angular difference between the {223}<692> orientation and a {111} plane is 11.4° and direction cosines are 0.98. As described previously, in the {111} plane, the [100] orientation being the orientation of easy magnetization is not contained, and orientation of this plane in the metal sheet plane is disadvantageous for obtaining a high magnetic flux density. Thus, the intensity C is set not to have the degree of contribution to improving a magnetic flux density, is put in the Z value as a parameter that performs division, and is multiplied by 0.98 being the direction cosines with respect to the {111} plane as its weighting.

From the above thought, it was found that when the intensity ratios in the respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in the metal sheet plane by X-ray diffraction are set to A, B, and C respectively, the expression of Z=(A+0.97B)/0.98C is created, and as the Z value is increased, a high magnetic flux density can be obtained when excitation is performed in the metal sheet plane to perform magnetometry.

Further, the present inventors were able to find from a large number of experiments that a special condition capable of obtaining a high magnetic flux density in an arbitrary direction in the metal sheet plane is that the Z value is not less than 2.0 nor more than 200. They grasped the fact that the Z value is limited to this range, and thereby the [100] orientation being the orientation of easy magnetization is thoroughly distributed in the metal sheet plane, but have not obtained evidence making theoretical explanation of this phenomenon possible so far.

The present inventors found that when the Z value is not less than 2.0 nor more than 200, a ratio B50/Bs of an average magnetic flux density B50 to a saturation magnetic flux density Bs becomes a high level of 0.80 or more and a magnetic flux density difference ΔB measured in the metal sheet plane becomes a low level of 0.15 T or less. FIG. 2 schematically shows this relationship.

When the Z value is less than 2.0, crystal orientation of α-Fe shows a tendency to decrease the [100] orientations being the orientation of easy magnetization in the metal sheet plane. Alternately, it shows a tendency that the distribution of the [100] orientations in the metal sheet plane becomes non-uniform. That is, the average magnetic flux density B50 in the metal sheet plane becomes small and the ratio B50/Bs of the average magnetic flux density B50 to the saturation magnetic flux density Bs becomes less than 0.8. Alternately, only the magnetic flux density in a specific direction increases and the magnetic flux density difference AB becomes greater than 0.15 T. Thus, the Z value is set to 2.0 or more in the present invention.

On the other hand, when the Z value exceeds 200, the increase in the magnetic flux density is saturated and an increase in uniformity of the magnetic flux density in the metal sheet plane is also saturated. In contrast to this, in order to manufacture a metal sheet such that the Z value exceeds 200, a heat treatment time is prolonged, or the like, which becomes difficult industrially, and thus the condition of the Z value is set to 200 or less.

Here, FIG. 1 is a view for explaining a method of calculating the average magnetic flux density B50. A manufacturing method will be described later, but it is found that α-region rolling is performed at 800° C. and as a different metal, 2.6 mass % of Sn and 0.9 mass % of Al are used, and thereby in an obtainable Fe-based metal sheet having a thickness of 0.2 mm, a high magnetic flux density can be obtained thoroughly in an in-plane circumferential direction.

Here, in a metal sheet having a higher accumulation degree of {200} planes among textures of the Fe-based metal sheet of the present invention in which the Z value is not less than 2.0 nor more than 200, a higher magnetic flux density can be obtained. Specifically, in a texture in which an accumulation degree of {200} planes in an α-Fe phase is not less than 30% nor more than 99% and an accumulation degree of {222} planes in the α-Fe phase is not less than 001% nor more than 30%, a higher magnetic flux density can be obtained.

When the accumulation degree of the {200} planes is less than 30% or the accumulation degree of the {222} planes is greater than 30%, the average magnetic flux density B50 tends to slightly decrease even though the Z value is in the present invention range. Further, in a metal sheet in which the accumulation degree of the {200} planes is greater than 99% or the accumulation degree of the {222} planes is less than 0.01%, the increase in the magnetic flux density B50 is saturated and a heat treatment time is prolonged, and the like, resulting in that manufacturing conditions become disadvantageous industrially.

Next, the manufacturing method of the previously described Fe-based metal sheet will be explained.

First Embodiment

As a manufacturing method of an Fe-based metal sheet in this embodiment, a rolling temperature and a reduction ratio are optimized, and thereby a {100} texture is formed in at least a surface layer portion of the metal sheet, a ferrite-forming element is made to diffuse into this partial or whole region from its surface, and at the time of cooling, the whole Fe-based metal sheet is oriented in {100}. This makes it possible to obtain a high magnetic flux density in an arbitrary direction in a metal sheet plane.

This embodiment as above is based on the fact found by the present inventors that {100} crystal grains in the texture formed in the surface layer portion preferentially grow at an A3 point or higher in a heating process to be performed for the diffusion of the ferrite-forming element, and further when the ferrite-forming element is made to diffuse into the inner portion to make the Fe-based metal sheet alloyed therewith and then cooling is performed, an accumulation degree of {200} planes in the sheet plane of the Fe-based metal sheet increases.

[Explanation of the Basic Principle of the First Embodiment of the Present Invention]

First, the basic principle of this embodiment capable of obtaining a high accumulation degree of {200} planes will be explained based on FIG. 3A to FIG. 3E.

(a) Manufacture of a Base Metal Sheet (seeding of a Texture)

In a process in which a cast slab containing C: less than 0.02 mass % and made of an Fe-based metal of an α-γ transforming component is decreased in thickness by rolling and thereby a metal sheet is obtained, hot rolling is performed at a sheet temperature of the A3 point or higher, α-region rolling is performed at a sheet temperature of lower than the A3 point and 300° C. or higher, and further cold rolling is performed to a predetermined sheet thickness. By this process, as shown in FIG. 3A, a base metal sheet 1 having an inner region 4 made of Fe in an a phase and having a {100} texture 2 in at least a surface layer portion 3 is obtained. Further, a seed of crystal that satisfies the condition of the Z value is formed in a recrystallized texture by a particular deformation slip.

(b) (Formation of a Second Layer)

Next, as shown in FIG. 3B, the ferrite-forming element such as Al, for example, is bonded to one surface or both surfaces of the cold-rolled base metal sheet 1 by using a vapor deposition method or the like to form a second layer 5.

(c) Saving of the Texture

Next, the base metal sheet 1 having had the ferrite-forming element bonded thereto is heated to the A3 point of the base metal sheet 1 to make the ferrite-forming element diffuse into the partial or whole region having the {100} texture 2 in the base metal sheet 1, to make the base metal sheet 1 alloyed therewith. As shown in FIG. 3C, an alloyed region 6 is transformed to the α phase from a γ phase to have an α single phase component. At this time, the alloyed region 6 is transformed while taking over orientation of the {100} texture 2 formed in the surface layer portion 3, so that a structure oriented in {100} is formed also in the alloyed region 6.

(d) Achievement of High Accumulation of the Texture

Next, the partially alloyed base metal sheet 1 is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region of the α single phase component is an α-Fe phase not undergoing γ transformation, and thus the {100} crystal grains are maintained as they are, the {100} crystal grains preferentially grow in the region, and the accumulation degree of the {200} planes increases. Further, as shown in FIG. 3D, a region 8 not having the α single phase component is transformed to the γ phase from the α phase.

Further, when a holding time of the temperature after the heating is prolonged, the {100} crystal grains are united to preferentially grow to large {100} crystal grains 7. As a result, the accumulation degree of the {200} planes further increases. Further, with the diffusion of the ferrite-forming element, the region 6 alloyed with the ferrite-forming element is transformed to the α phase from the γ phase. At this time, in the region adjacent to the region to be transformed, crystal grains in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the region 6 is transformed while taking over a crystal orientation of the adjacent crystal grains in the α phase. Thereby, the holding time is prolonged and the accumulation degree of the {200} planes increases.

(e) Growth of the Texture

The base metal sheet is cooled to a temperature of lower than the A3 point. At this time, as shown in FIG. 3E, a γ-Fe phase in an unalloyed inner region 10 is transformed to the α-Fe phase. This inner region 10 is adjacent to the region in which the crystal grains in the α phase oriented in {100} are already formed in a temperature region of the A3 point or higher, and at the time of the transformation to the α phase from the γ phase, the inner region 10 is transformed while taking over the crystal orientation of the adjacent crystal grains in the α phase and larger crystal grains 9 in the α phase oriented in {100} are formed. Therefore, the accumulation degree of the {200} planes increases also in the region. By this phenomenon, the high accumulation degree of the {200} planes can be obtained even in the unalloyed region.

When at the stage of the preceding state shown in FIG. 3D, the temperature of the A3 point or higher is held until the whole metal sheet is alloyed, the structure having the high accumulation degree of the {200} planes is already formed in the whole metal sheet, and thus the cooling is performed while the state when the cooling is started is maintained.

In the above, the basic principle of this embodiment was explained, and there will be further explained a limiting reason of each condition that defines the manufacturing method of this embodiment and preferable conditions of this embodiment.

[Fe-Based Metal to be the Base Material] (C Content)

In this embodiment, first, crystal grains oriented in {100} to serve as seeds for increasing the accumulation degree of the {200} planes in the sheet are formed in the surface layer portion of the base metal sheet made of the Fe-based metal. Then, the γ-α transformation is made to progress in the metal sheet while taking over a crystal orientation of the crystal grains in the α phase to serve as the seeds finally, to thereby increase the accumulation degree of the {200} planes of the whole metal sheet. For this reason, the Fe-based metal used for the base metal sheet has a composition of the α-γ transforming component. When the Fe-based metal used for the base metal sheet has the α-γ transforming component, the ferrite-forming element is made to diffuse into the metal sheet to make the metal sheet alloyed therewith, thereby making it possible to form the region having the α single phase based component.

In this embodiment, the C content of the base metal sheet is set to less than 0.02 mass %. Further, in terms of a magnetic property of a product metal sheet, the C content is preferably 0.01 mass % or less. Under the condition of the C content being less than 0.02 mass %, the ferrite-forming element is made to diffuse into the metal sheet to make the metal sheet alloyed therewith, thereby making it possible to form the region having the α single phase based component. Incidentally, C is a component to remain in a process of manufacturing the slab and the less C is, the more preferred it is in terms of the magnetic property, and thus its lower limit is not necessary needed, but it is preferably set to 0.0001 mass % or more in terms of the cost of a refining process.

(Other Containing Elements)

In principle, being applicable to the Fe-based metal having the α-γ transforming component, this embodiment is not limited to the Fe-based metal in a specific composition range. Typical examples of the α-γ transforming component are pure iron, steel such as ordinary steel, and the like. For example, it is a component containing pure iron or steel containing C of 1 ppm to less than 0.02 mass % as described above and a balance being composed of Fe and inevitable impurities as its base and containing an additive element as required. Instead, it may be silicon steel of the α-γ transforming component having C: less than 0.02 mass % and Si: 0.1 mass % to 2.5 mass % as its basic component. Further, as other impurities, a trace amount of Ni, Cr, Al, Mo, W, V, Ti, Nb, B, Cu, Zr, Y, Hf, La, Ce, N, O, P, S, and/or the like are/is contained. Further, Al and Mn are added to increase electric resistance, to thereby decrease a core loss, and Co is added to increase the saturation magnetic flux density Bs, to thereby increase a magnetic flux density, which are also included in the present invention range.

(Thickness of the Base Metal Sheet)

The thickness of the base metal sheet is set to not less than 10 μm nor more than 6 mm. When the thickness is less than 10 μm, when the base metal sheets are stacked to be used as a magnetic core, the number of the sheets to be staked is increased to increase gaps, resulting in that a high magnetic flux density cannot be obtained. Further, when the thickness exceeds 6 mm, it is not possible to make the {100} texture grow sufficiently even though a reduction ratio of the α-region rolling is adjusted, resulting in that a high magnetic flux density cannot be obtained.

[Rolling Conditions]

In this embodiment, as described previously, the Fe-based metal having, in at least the surface layer portion, the crystal grains oriented in {100} to serve as the seeds for increasing the accumulation degree of the {200} planes in the metal sheet is used as a starting material. As a method of achieving high accumulation of the {100} planes of the base metal sheet, a method of performing α-region rolling in a process in which a cast slab is rolled to a sheet shape is used.

First, a cast slab containing C: less than 0.02 mass % and made of the Fe-based metal of the α-γ transforming component such as a continuous cast slab or an ingot is prepared. Then, in a process in which the cast slab is decreased in thickness by rolling to obtain the base metal sheet, first the hot rolling is performed at a temperature of the A3 point or higher. Next, the α-region rolling is performed at a temperature of lower than the A3 point and higher than 300° C., and further the metal sheet is subjected to cold rolling to a predetermined thickness, and thereby the base metal sheet having the {100} texture formed in the surface layer portion is obtained.

As for a reduction ratio in each of rolling processes to be performed until the base metal sheet is obtained from the cast slab, the total reduction ratio in the α-region rolling is preferably set to −1.0 or less in terms of true strain and the sum of the total reduction ratio in the α-region rolling and the total reduction ratio in the cold rolling is preferably set to −2.5 or less in terms of true strain. Conditions other than these may create a possibility that the {100} texture cannot be sufficiently formed in the surface layer portion. A method of expressing the reduction ratio by true strain ε is expressed by the following expression (1), where in each of the rolling processes, the thickness before the rolling is set to h0 and the thickness after the rolling is set to h.

ε=ln(h/h0)   (1)

When the sum of the total reduction ratio in the α-region rolling and the total reduction ratio in the cold rolling is in the previously described preferred range, a deformed structure in which the {100} texture is formed by recrystallization can be provided to at least the vicinity of the surface layer portion of the base metal sheet. Particular crystal slip and crystal rotation to occur at these reduction ratios are thought to occur. Thus, they are preferably in these ranges.

Further, as for the reduction ratio in each of the rolling processes to be performed until the base metal sheet is obtained from the cast slab, the reduction ratio in the hot rolling is preferably −0.5 or less in terms of true strain, thereby making it easier to obtain the higher accumulation degree of the {200} planes. This results from the fact found by the present inventors that in order that desirable deformation should be performed in the α-region rolling and the cold rolling, deformation in the hot rolling in a γ region is also closely affected. Thus, these ranges are preferred.

The region of the surface layer portion in which the {100} texture is formed preferably has 1 μm or more of a distance in a sheet thickness direction from the surface. Thereby, it is possible to bring the accumulation degree of the {200} planes to 30% or more in the following diffusion treatment. The upper limit of the distance is not limited in particular, but it is difficult to form the {100} texture in a region of 500 μm or more by rolling.

Incidentally, the measurement of the accumulation degree of the {200} planes can be performed by X-ray diffraction using a MoKα ray. To be in more detail, in the α-Fe crystal, integrated intensities of 11 orientation planes ({110}, {200}, {211}, {310}, {222}, {321}, {411}, {420}, {332}, {521}, and {442}) parallel to a sample surface are measured for each sample, each measured value is divided by a theoretical integrated intensity of the sample having a random orientation, and thereafter, a ratio of the intensity of {200} or {222} is obtained in percentage.

At this time, for example, the accumulation degree of the {200} planes is expressed by Expression (2) below.

accumulation degree of {200} planes=[{i(200)/I(200)}/Σ{i(hkl)/I(hkl)}]'100   (2)

Here, i(hkl) is an actually measured integrated intensity of {hkl} planes in the measured sample, and I(hkl) is a theoretical integrated intensity of the {hkl} planes in the sample having the random orientation. Further, Σ is the sum of the 11 orientation planes in the α-Fe crystal. Here, instead of the theoretical integrated intensity of the sample having the random orientation, actually measured values using the sample may be used.

[Different Metal]

Next, a different metal except Fe is made to diffuse into the base metal sheet manufactured by the above-described rolling processes to increase the region of the {100} texture in the thickness direction of the steel sheet. As the different metal, the ferrite-forming element is used. As a procedure, first, the different metal is bonded in a layered form as the second layer to one surface or both surfaces of the base metal sheet made of the Fe-based metal of the α-γ transforming component. Then, a region alloyed by having had elements of the different metal diffuse thereinto is turned to have the α single phase based component and to be able to be maintained as not only the region transformed to the α phase, but also a seed oriented in {100} for increasing the accumulation degree of the {200} planes in the metal sheet. As such a ferrite-forming element, at least one type of Al, Cr, Ga, Ge, Mo, Sb, Si, Sn, Ta, Ti, V, W, and Zn can be used alone or in a combined manner.

As a method of bonding the different metal in a layered form to the surface of the base metal sheet, there can be employed various methods such as a plating method of hot dipping, electrolytic plating, or the like, a rolling clad method, a dry process of PVD, CVD, or the like, and further powder coating. As a method of efficiently bonding the different metal for industrially implementing the method, the plating method or the rolling clad method is suitable.

The thickness of the different metal before the heating when the different metal is bonded is preferably not less than 0.05 μm nor more than 1000 μm. When the thickness is less than 0.05 μm, it is not possible to obtain the sufficient accumulation degree of the {200} planes. Further, when the thickness exceeds 1000 μm, even when the different metal layer is made to remain, its thickness becomes larger than necessary.

[Heating and Diffusion Treatment]

The base metal sheet having had the ferrite-forming element as the different metal bonded thereto is heated up to the A3 point of the base metal sheet, to thereby make the ferrite-forming element diffuse into the partial or whole region of the {100} texture formed in the surface layer portion of the base metal sheet to make the base metal sheet alloyed therewith. The region alloyed with the ferrite-forming element is turned to have the α single phase component and the region is transformed to the α phase from the γ phase. At this time, the region is transformed while taking over the orientation of the {100} texture formed in the surface layer portion, and thus the structure oriented in {100} is formed also in the alloyed region. As a result, in the alloyed region, a structure in which the accumulation degree of the {200} planes in the α-Fe phase becomes not less than 25% nor more than 50% and in accordance with it, the accumulation degree of the {222} planes in the α-Fe phase becomes not less than 1% nor more than 40% is formed.

Then, the base metal sheet is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region alloyed already is turned into an α single phase structure that is not transformed to the γ phase, so that the crystal grains in the {100} texture are maintained as they are, and in the region, the crystal grains in the {100} texture preferentially grow and the accumulation degree of the {200} planes increases. Further, the region not having the α single phase component is transformed to the γ phase.

Further, when the holding time is prolonged, the crystal grains in the {100} texture are united to one another to preferentially grow. As a result, the accumulation degree of the {200} planes further increases. Further, with the further diffusion of the ferrite-forming element, the region alloyed with the ferrite-forming element is transformed to the α phase from the γ phase. At this time, as shown in FIG. 4A, in the regions adjacent to the regions to be transformed, crystal grains 7 in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the regions alloyed with the ferrite-forming element are transformed while taking over a crystal orientation of the adjacent crystal grains 7 in the α phase. By these phenomena, the holding time is prolonged and the accumulation degree of the {200} planes increases. Further, as a result, the accumulation degree of the {222} planes decreases.

Incidentally, in order to finally obtain the high accumulation degree of the {200} planes of 50% or more, it is preferred that the holding time should be adjusted to, at this stage, bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more and bring the accumulation degree of the {222} planes in the α-Fe phase to 30% or less. Further, when the A3 point or higher is held until the whole metal sheet is alloyed, as shown in FIG. 4C, the α single phase structures are formed up to the center portion of the metal sheet and grain structures oriented in {100} reach the center of the metal sheet.

A holding temperature after the temperature is increased is set to not lower than the A3 point nor higher than 1300° C. Even when the metal sheet is heated at a temperature higher than 1300° C., an effect with respect to the magnetic property is saturated. Further, cooling may be started immediately after the temperature reaches the holding temperature (in the case, the temperature is held for 0.01 second or longer substantially), or cooling may also be started after the temperature is held for 600 minutes or shorter. Even when the temperature is held for longer than 600 minutes, the effect is saturated. When this condition is satisfied, the achievement of high accumulation of the seeds oriented in the {200} plane further progresses to make it possible to more securely bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more after the cooling.

[Cooling After the Heating and Diffusion Treatment]

After the diffusion treatment, when the cooling is performed while the region that is not alloyed with the ferrite-forming element is remaining, as shown in FIG. 4B, at the time of the transformation to the α phase from the γ phase, the unalloyed region is transformed while taking over the crystal orientation of the regions in which the crystal grains 9 in the α phase oriented in {100} are already formed. Thereby, the accumulation degree of the {200} planes increases, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {222} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained, the crystal satisfying the condition of the Z value grows, and a high magnetic flux density can be obtained in an arbitrary direction in the metal sheet plane.

Further, as shown in FIG. 4C, when the A3 point or higher is held until the whole metal sheet is alloyed, and the grain structures oriented in {100} reach the center of the metal sheet, as shown in FIG. 4D, the metal sheet is cooled as it is, and the texture in which the crystal grains 9 oriented in {100} reach up to the center of the metal sheet can be obtained. Thereby, the whole metal sheet is alloyed with the different metal, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {222} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained.

As above, the value of the accumulation degree of the {200} planes and the remaining state of the different metal on the surface of the base metal sheet change depending on the holding time of the temperature of the A3 point or higher and the holding temperature. The example shown in FIG. 4B is in a state where the grain structures oriented in {100} do not reach up to the center of the metal sheet, the different metal also remain on the surfaces, and an α single phase front surface side region and an α single phase rear surface side region being the second layer are formed, but it is also possible to obtain the grain structures oriented in {100} up to the center of the metal sheet and to alloy all the second layers on the surfaces.

Incidentally, at the time of the cooling after the diffusion treatment, a cooling rate is preferably not less than 0.1° C./sec nor more than 500° C./sec. When the cooling is performed in this temperature range, the growth of the seeds oriented in the {200} plane further progresses.

Further, when the second layers are made to remain on the obtainable Fe-based metal sheet having a thickness of not less than 10 μm nor more than 6 mm, the thickness of the second layer is preferably set to not less than 0.01 μm nor more than 500 μm. Further, a ratio of the α single phase region alloyed at this stage is preferably 1% or more in a cross section of the Fe-based metal sheet.

Further, at the time of cooling to a temperature of lower than the A3 point in the state shown in FIG. 4A, an average cooling rate at the time of cooling to the A3 point −50° C. from the A3 point may be set to 50° C./minute or less. When the cooling is performed at the cooling rate in this range, the adjacent crystal grains oriented in {100} are united to one another to grow, and as shown in FIG. 5, a coarse crystal grain 11 straddling part of an α single phase front surface side region 6 a adjacent to a front surface side second layer 5 a and straddling part of an α single phase rear surface side region 6 b adjacent to a rear surface side second layer 5 b is formed. When the average cooling rate from the A3 point to the A3 point −50° C. becomes greater than 50° C./minute, there is no sufficient time for growth of the crystal grain 11, resulting in that an excellent core loss property cannot be obtained. On the other hand, the lower limit of the average cooling rate from the A3 point to the A3 point −50° C. is not limited, but the lower limit is preferably set to 1° C./minute in terms of the productivity.

Further, in order to obtain a more excellent core loss property, an average cooling rate at the time of cooling to the A3 point −10° C. from the A3 point is preferably set to 20° C./minute or less. On the other hand, the lower limit of the average cooling rate from the A3 point to the A3 point −10° C. is not limited, but the lower limit is preferably set to 1° C./minute in terms of the productivity.

Second Embodiment

In the previously described first embodiment, there was explained the manufacturing method of the previously described Fe-based metal sheet by using the cast slab containing C: less than 0.02 mass % and made of the Fe-based metal of the α-γ transforming component. In contrast to this, in this embodiment, there will be explained a manufacturing method of the previously described Fe-based metal sheet by using a cast slab containing C: 0.02 mass % or more.

When the C content is large, a good magnetic property cannot be obtained, so that it is necessary to remove C by performing decarburization annealing. Thus, the decarburization annealing is performed under conditions to be explained below, thereby making it possible to increase the accumulation degree of the {200} planes.

In the method of this embodiment, a {100} texture is formed in a surface layer portion of an Fe-based metal sheet by using γ-α transformation accompanying decarburization (and further demanganization), and thereafter a ferrite-forming element is made to diffuse into a partial or whole decarburized region and further over the region from its surface, and at the time of cooling, the whole Fe-based metal sheet is made to be oriented in {100}.

This embodiment as above is based on the fact found by the present inventors that {100} crystal grains in the texture formed in the surface preferentially grow at an A3 point or higher in a heating process to be performed for the diffusion of the ferrite-forming element and further when the ferrite-forming element is made to diffuse into the inner portion to make the Fe-based metal sheet alloyed therewith and then cooling is performed, an accumulation degree of {200} planes in a sheet plane of the Fe-based metal sheet increases.

[Explanation of the Basic Principle of the Second Embodiment of the Present Invention]

First, the basic principle of this embodiment capable of obtaining a high accumulation degree of {200} planes will be explained based on FIG. 3A to FIG. 3D, by taking the case of decarburization as an example.

(a) Seeding of a Texture

When being decarburized until C becomes less than 0.02 mass %, the Fe-based metal sheet containing C: 0.02 mass % or more and having a composition of the α-γ transforming component is heated to a temperature at which a structure is turned into an α single phase and to a temperature of a γ single phase or a two-phase region of a γ phase and an α phase (namely, a temperature of an A1 point or higher) to decarburize the surface layer portion of the Fe-based metal sheet until C becomes less than 0.02 mass %. Thereby, the γ-α transformation is made to occur in a process of the decarburization to turn the decarburized surface layer portion into the α phase.

At this time, the decarburization progresses the most in the <110> direction in the γ phase having large gaps between lattices, and in this portion, the C concentration becomes less than 0.02% and the transformation to the α phase occurs. A {110} plane in the γ phase becomes the {100} plane when the γ phase is turned into the α phase in a BCC structure, and thus in the α phase after the decarburization, the {100} plane is preferentially formed. Further, the growth, of the crystal grains in the α phase formed in the surface, in the sheet thickness direction is slow because its rate is controlled by a decarburization rate, and thus the crystal grains in the α phase formed in the surface grow in a direction parallel to the sheet plane. Further, in the surface of the metal sheet, the {100} plane preferentially grows by taking surface energy as driving force. As a result, the whole surface of the metal sheet becomes a structure oriented in {100} finally. By this process, as shown in FIG. 3A, a base metal sheet 1 having an inner region 4 made of Fe in the α phase and having the accumulation degree of the {200} planes in the decarburized region brought to 20% or more can be obtained. Further, a seed of crystal that satisfies the condition of the Z value is formed in the structure formed at the time of the decarburization by taking the surface energy as driving force.

(b) (Formation of a Second Layer)

Next, as shown in FIG. 3B, the ferrite-forming element such as Al is bonded to one surface or both surfaces of the base metal sheet 1 after the decarburization by using a vapor deposition method or the like to form a second layer 5.

(c) Saving of the Texture

Next, the base metal sheet 1 having had the ferrite-forming element bonded thereto is heated to the A3 point of the base metal sheet 1 to make the ferrite-forming element diffuse into the partial or whole decarburized region in the base metal sheet 1 to make the base metal sheet 1 alloyed therewith. Thereby, as shown in FIG. 3C, the α phase is formed in an alloyed region 6. Alternately, the ferrite-forming element is made to diffuse into the inner portion over the decarburized region to make the base metal sheet 1 alloyed therewith, and the alloyed region is turned to have the α single phase component partially, and thereby the region is turned into the α phase. At this time, the region is transformed while taking over orientation of the region formed by the decarburization, so that the structure oriented in {100} is formed also in the alloyed region 6. Further, the orientation in {100} is further increased even in crystal grains turned into the α phase previously. Further, when the ferrite-forming element is made to diffuse and the crystal is oriented, the seed of the crystal satisfying the condition of the Z value preferentially grows.

(d) Achievement of High Accumulation of the Texture

Next, the partially alloyed base metal sheet 1 is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region of the α single phase component is an α-Fe phase not undergoing γ transformation, and thus the {100} crystal grains are maintained as they are, the {100} crystal grains preferentially grow in the region, and the accumulation degree of the {200} planes increases. Further, as shown in FIG. 3D, the region not having the α single phase component is transformed to the γ phase from the α phase.

Further, when a holding time of the temperature after the heating is prolonged, the {100} crystal grains are united to preferentially grow to large {100} crystal grains 7. As a result, the accumulation degree of the {200} planes further increases. Further, with the diffusion of Al, the region alloyed with Al is transformed to the α phase from the γ phase. At that time, in the region adjacent to the region to be transformed, crystal grains in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the region is transformed while taking over a crystal orientation of the adjacent crystal grains in the α phase. Thereby, the holding time is prolonged and the accumulation degree of the {200} planes increases.

(e) Growth of the Texture

Next, the base metal sheet is cooled to a temperature of lower than the A3 point. At this time, as shown in FIG. 3E, a γ-Fe phase in an unalloyed inner region 10 is transformed to the α-Fe phase. This inner region 10 is adjacent to the region in which the crystal grains in the α phase oriented in {100} are already formed in a temperature region of the A3 point or higher, and at the time of the transformation to the α phase from the γ phase, the inner region 10 is transformed while taking over the crystal orientation of the adjacent crystal grains in the α phase and larger crystal grains 9 in the α phase oriented in {100} are formed. Therefore, the accumulation degree of the {200} planes increases also in the region (see the state shown in FIG. 3E). By this phenomenon, the high accumulation degree of the {200} planes can be obtained even in the unalloyed region 10.

When at the stage of the preceding state shown in FIG. 3D, the temperature of the A3 point or higher is held until the whole metal sheet is alloyed, the structure having the high accumulation degree of the {200} planes is already formed in the whole metal sheet, and thus the cooling is performed while the state when the cooling is started is maintained.

Further, in the above explained example, the Fe-based metal sheet containing C: 0.02 mass % or more is used, but when an Fe-based metal sheet containing C: less than 0.02 mass % is used, carburization is performed before the decarburization to bring the C content in the region to be decarburized to 0.02 mass % or more.

In the above, the basic constitution of this embodiment was explained, and there will be further explained a limiting reason of each condition that defines a manufacturing method of this embodiment and preferable conditions of this embodiment.

[Fe-Based Metal to be the Base Material] (C Content)

In this embodiment, first, crystal grains oriented in {100} to serve as seeds for increasing the accumulation degree of the {200} planes are formed in the surface of the base metal sheet made of the Fe-based metal. Then, the γ-α transformation is made to progress in the metal sheet while taking over a crystal orientation of the crystal grains in the α phase to serve as the seeds finally, to thereby increase the accumulation degree of the {200} planes of the whole metal sheet.

In this embodiment, the seeds of the crystal grains oriented in {100} are formed in the surface of the base metal sheet by structure control using the γ-α transformation accompanying decarburization or demanganization. The Fe-based metal used for the base metal sheet has a composition of the α-γ transforming component, and the C content in the region to be decarburized is brought to 0.02 mass % or more.

Further, the Fe-based metal used for the base metal sheet has the α-γ transforming component, and the ferrite-forming element is made to diffuse into the metal sheet to make the metal sheet alloyed therewith, thereby making it possible to form a region having the α single phase based component. Further, the C content in the region to be decarburized is brought to 0.02 mass % or more, thereby making it possible to use the γ-α transformation accompanying the decarburization.

For bringing the C content in the base metal sheet to 0.02 mass % or more, there is a method of using a base metal sheet manufactured from a molten material adjusted to contain C: 0.02 mass % or more by undergoing casting and rolling processes (a melting method). As another method, there is a method in which a base metal sheet having the C content of less than 0.02 mass % is used and in a surface layer portion of the base metal sheet, a region containing C: 0.02 mass % or more is formed by carburization.

In the case of the melting method, the range of the C content is set to not less than 0.02 mass % nor more than 1.0 mass %. When the C content is less than 0.02 mass %, it is not possible to use the formation of a {200} texture using the γ-α transformation accompanying the decarburization. Further, when the C content is more than 1.0% mass, a long time is required for the decarburization. The preferred range of the C content is not less than 0.05 mass % nor more than 0.5 mass %.

In the case of the carburization method, the range of the C content of the Fe-based metal of which the base metal sheet is made is set to 1 ppm or more and less than 0.02 mass %. Then, the surface layer of this Fe-based metal is subjected to the carburization so that the C concentration may become not less than 0.02 mass % nor more than 1.0 mass % in the same manner as that in the melting case.

Further, a carburizing range is set to a region down to a distance y from the surface, where the distance in a depth direction from the surface is set to y. This distance y is not less than 5 μm nor more than 50 μm. When the distance y is less than 5 μm, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment after the decarburization, so that the distance y is set to 5 μm or more. Further, when the distance becomes greater than 50 μm, a long time is required for the carburization, and further a long time is required also for the decarburization of the whole carburized region. Further, an obtainable effect is also saturated, so that the preferred distance y is set to 50 μm or less. The carburizing method is not limited in particular, and a well-known gas carburizing method or the like may be performed.

Incidentally, the C content is preferably 0.005 mass % or less in terms of a magnetic property of a product metal sheet, so that in order to manufacture a steel sheet excellent in a magnetic property, silicon steel having the C content of 0.005 mass % or less is used to be subjected to carburization in a manner to have the above-described C concentration, which is advantageous for cost.

(Mn Content)

When Mn being an austenite stabilizing element is contained in the Fe-based metal, it is possible to form seeds of crystal grains oriented in {100} by structure control using the γ-α transformation accompanying demanganization. The demanganization is performed together with the decarburization, and thereby the surface layer portion is turned into the α phase more efficiently and the accumulation degree of the {200} planes in a decarburized and demanganized region is more increased. In order to exhibit such a function, the Mn content before performing the demanganization treatment is preferably set to 0.2 mass % or more.

The above-described structure control using the γ-α transformation can be performed even by the decarburization alone, so that Mn does not have to be contained. However, when Mn is contained, an effect of increasing electric resistance to decrease a core loss is also obtained, and thus Mn in a range of 2.0 mass % or less may also be contained according to need even when no demanganization is performed. From the above point, the range of the Mn content when Mn is contained is preferably set to 0.2 mass % to 2.0 mass %.

(Other Containing Elements)

In principle, being applicable to the Fe-based metal having the α-γ transforming component, this embodiment is not limited to the Fe-based metal in a specific composition range. Typical examples of the α-γ transforming component are pure iron, steel such as ordinary steel, and the like. For example, it is a component containing pure iron or steel containing C of 1 ppm to 0.10 mass % as described above or further containing Mn of 0.2 mass % to 2.0 mass % and a balance being composed of Fe and inevitable impurities as its base and containing an additive element as required. Instead, it may be silicon steel of the α-γ transforming component having C: 1.0 mass % or less and Si: 0.1 mass % to 2.5 mass % as its basic component. Further, as other impurities, a trace amount of Ni, Cr, Al, Mo, W, V, Ti, Nb, B, Cu, Co, Zr, Y, Hf, La, Ce, N, O, P, S, and/or the like are/is contained. Incidentally, Al and Mn are added to increase electric resistance, to thereby decrease a core loss and Co is added to increase the saturation magnetic flux density Bs, to thereby increase a magnetic flux density, which are also included in the present invention range.

(Thickness of the Base Metal Sheet)

The thickness of the base metal sheet is set to not less than 10 μm nor more than 6 mm. When the thickness is less than 10 μm, when the base metal sheets are stacked to be used as a magnetic core, the number of the sheets to be staked is increased to increase gaps, resulting in that a high magnetic flux density cannot be obtained. Further, when the thickness is greater than 6 mm, it is not possible to make the {100} texture grow sufficiently after cooling after the diffusion treatment, resulting in that a high magnetic flux density cannot be obtained.

[Decarburization Treatment]

In the decarburization treatment for turning the surface layer portion of the base metal sheet into the α phase, the base metal sheet is desirably heated in a decarburizing atmosphere to be decarburized in the following manner.

(Temperature of the Decarburization Treatment)

The temperature of the decarburization treatment is set to a temperature of the A1 point or higher and a temperature at which a structure is turned into an α single phase when the decarburization is performed until C becomes less than 0.02 mass %. The base metal sheet containing C: 0.02 mass % or more is heated to a temperature of a γ single phase or a two-phase region of a γ phase and an α phase (namely a temperature of the A1 point or higher) in order to make the γ-α transformation occur by the decarburization.

(Atmosphere of the Decarburization Treatment)

With regard to the decarburizing atmosphere, a conventionally known method in manufacture of a grain-oriented electrical steel sheet can be employed. For example, there is a method in which decarburization is first performed in a weak decarburizing atmosphere, in a vacuum of 1 Torr or less, for example, or in a gas atmosphere of one type or two or more types of H₂, He, Ne, Nr, Kr, Xe, Rn, and N₂ at a temperature of lower than (a dew point −20)° C., and next decarburization is performed in a strong decarburizing atmosphere, or in a gas atmosphere in which an inert gas, or CO and CO₂ is/are added to H₂ at a temperature of (a dew point −20)° C. or higher, for example. In this case, if the decarburization is continued to the end in the weak decarburizing atmosphere, a long time is required.

(Period of Performing the Decarburization Treatment)

The period of performing the decarburization treatment is preferably not shorter than 0.1 minute nor longer than 600 minutes. When the period is shorter than 0.1 minute, it is difficult to bring the accumulation degree of the {200} planes to 20% or more after the decarburization, and when the period is long so as to exceed 600 minutes, too much cost is needed.

(Range of Performing the Decarburization Treatment)

The range of performing the decarburization treatment is a range down to a distance x, where the distance in the depth direction from the surface is set to x, and the distance x is not less than 5 μm nor more than 50 μm. When the distance x is less than 5 μm, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment after the decarburization. For this reason, the distance x in the depth direction from the surface is set to 5 μm or more. Further, when the distance is greater than 50 μm, a long time is required for the decarburization, and further the accumulation degree of the {200} planes is saturated, and thus it is not advantageous industrially. Thus, the distance x is set to 50 μm or less.

(Other Decarburizing Methods)

Further, as described in Patent Literature 6, it is also possible that a material promoting decarburization is applied to a surface of a steel sheet as an annealing separating agent and this is wound around a coil and is subjected to coil annealing, to thereby form a decarburized region. Further, it is also possible that the above-described annealing separating agent is applied to a surface of a steel sheet in a single sheet form and the steel sheets are stacked to be subjected to annealing at the above-described temperature for a similar time, to thereby form a decarburized region.

(C Content after the Decarburization)

The C content after the decarburization is set to less than 0.02 mass % in order to obtain an α-phase single phase structure as described above. It is preferably 0.005 mass % or less in terms of the magnetic property of a product.

(Accumulation Degree of the {200} Planes after the Decarburization)

It is preferred that the accumulation degree of the {200} planes in the decarburized region after the decarburization should become 20% or more by performing the decarburization annealing under the above conditions. When the accumulation degree of the {200} planes is less than 20%, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment to be performed subsequently. Further, the upper limit of the accumulation degree of the {200} planes is preferably set to 99%. When it is greater than 99%, the magnetic property deteriorates. The accumulation degree of the {200} planes is adjusted to fall within the above-described range by selecting the conditions of the decarburizing temperature, the decarburizing time, the decarburizing atmosphere, and the like. Incidentally, the measurement of the accumulation degree of the plane in the above-described orientation plane can be performed by X-ray diffraction using a MoKα ray similarly to the first embodiment.

[Demanganization Treatment]

In this embodiment, the decarburization treatment and the demanganization treatment may also be used in combination by containing Mn in the base metal sheet. The demanganization treatment is performed simultaneously with the decarburization or subsequently to the decarburization under the following conditions. Incidentally, as described in Patent Literature 6, it is also possible to perform the decarburization treatment and the demanganization treatment simultaneously in a state where steel sheets each have an annealing separating agent containing a material promoting decarburization and a material promoting demanganization applied thereto to be staked.

(Temperature and Range of the Demanganization Treatment)

The temperature at which the demanganization treatment is performed is set to a temperature of the A1 point or higher similarly to the decarburization. With regard to a demanganizing atmosphere, the demanganization treatment may be performed under a reduced pressure atmosphere. Further, the period of performing the demanganization treatment is preferably set to fall within a range of not shorter than 0.1 minute nor longer than 600 minutes similarly to the decarburization.

(Range of Performing the Demanganization Treatment)

The range of performing the demanganization treatment is a range down to a distance x, where the distance in the depth direction from the surface is set to x, and the distance x is preferably not less than 5 μm nor more than 50 μm. When the distance x is less than 5 μm, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment after the demanganization. For this reason, the preferred distance x in the depth direction from the surface is set to 5 μm or more. Further, when the distance is greater than 50 μm, a long time is required for the demanganization, and further the accumulation degree of the {200} planes is saturated, and thus it is not advantageous industrially. Thus, the preferred distance x is set to 50 μm or less.

(Accumulation Degree of the {200} Planes after the Demanganization)

It is preferred that the accumulation degree of the {200} planes in the region having been subjected to the demanganization treatment should become 20% or more after the demanganization by performing the decarburization annealing under the above conditions. When the accumulation degree of the {200} planes is less than 20%, it is difficult to bring the accumulation degree of the {200} planes to 30% or more in the diffusion treatment to be performed subsequently. The upper limit of the accumulation degree of the {200} planes is preferably set to 99%. When it is greater than 99%, the magnetic property deteriorates.

[Different Metal]

Next, a different metal except Fe is made to diffuse into the base metal sheet having had the surface layer portion turned into the α phase by the decarburization to increase the region of the {100} texture in the thickness direction of the metal sheet. As the different metal, the ferrite-forming element is used. As a procedure, first, the different metal is bonded in a layered form as the second layer to one surface or both surfaces of the base metal sheet made of the Fe-based metal of the α-γ transforming component. Then, a region alloyed by having had elements of the different metal diffuse thereinto is turned to have the α single phase based component and to be able to be maintained as not only the region having been subjected to the decarburization (or further the demanganization) to be transformed to the α phase, but also a seed oriented in {100} for increasing the accumulation degree of the {200} planes in the metal sheet. As such a ferrite-forming element, at least one type of Al, Cr, Ga, Ge, Mo, Sb, Si, Sn, Ta, Ti, V, W, and Zn can be used alone or in a combined manner.

As a method of bonding the different metal in a layered form to the surface of the base metal sheet, there can be employed various methods such as a plating method of hot dipping, electrolytic plating, or the like, a rolling clad method, a dry process of PVD, CVD, or the like, and further powder coating. As a method of efficiently bonding the different metal for industrially implementing the method, the plating method or the rolling clad method is suitable.

The thickness of the different metal before the heating when the different metal is bonded is preferably not less than 0.05 μm nor more than 1000 μm. When the thickness is less than 0.05 μm, it is not possible to obtain the sufficient accumulation degree of the {200} planes. Further, when the thickness exceeds 1000 μm, even when the different metal layer is made to remain, its thickness becomes larger than necessary.

[Heating and Diffusion Treatment]

The base metal sheet having had the ferrite-forming element bonded thereto is heated up to the A3 point of the base metal sheet, to thereby make the ferrite-forming element diffuse into the partial or whole region in the base metal sheet to make the base metal sheet alloyed therewith. The α phase is maintained in the region alloyed with the ferrite-forming element. Alternately, the ferrite-forming element is made to diffuse into the inner portion over the decarburized region to make the base metal sheet alloyed therewith, and the alloyed region is turned to have the α single phase component partially, and thereby the region is turned into the α phase. At this time, the region is transformed while taking over the orientation of the region formed by the decarburization, so that the accumulation degree of the {200} planes further increases. As a result, in the alloyed region, a structure in which the accumulation degree of the {200} planes in the α-Fe phase becomes not less than 25% nor more than 50% and in accordance with it, the accumulation degree of the {222} planes in the α-Fe phase becomes not less than 1% nor more than 40% is formed.

Then, the base metal sheet is further heated to a temperature of not lower than the A3 point nor higher than 1300° C. and the temperature is held. The region alloyed already is turned into an α single phase structure that is not transformed to the γ phase, so that the {100} crystal grains are maintained as they are, and in the region, the crystal grains in the {100} texture preferentially grow and the accumulation degree of the {200} planes increases. Further, the region not having the α single phase component is transformed to the γ phase.

Further, when the holding time is prolonged, the crystal grains in the {100} texture are united to one another to preferentially grow. As a result, the accumulation degree of the {200} planes further increases. Further, with the further diffusion of the ferrite-forming element, the region alloyed with the ferrite-forming element is transformed to the α phase from the γ phase. At this time, as shown in FIG. 4A, in the regions adjacent to the regions to be transformed, crystal grains 7 in the α phase oriented in {100} are already formed, and at the time of the transformation to the α phase from the γ phase, the regions alloyed with the ferrite-forming element are transformed while taking over a crystal orientation of the adjacent crystal grains 7 in the α phase. Thereby, the holding time is prolonged and the accumulation degree of the {200} planes increases. Further, as a result, the accumulation degree of the {222} planes decreases.

Incidentally, in order to finally obtain the high accumulation degree of the {200} planes of 50% or more, it is preferred that the holding time should be adjusted to, at this stage, bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more and bring the accumulation degree of the {222} planes in the α-Fe phase to 30% or less. Further, when the A3 point or higher is held until the whole metal sheet is alloyed, as shown in FIG. 4C, the α single phase structures are formed up to the center portion of the metal sheet and grain structures oriented in {100} reach the center of the metal sheet.

A holding temperature after the temperature is increased is set to not lower than A3 point nor higher than 1300° C. Even when the metal sheet is heated at a temperature higher than 1300° C., an effect with respect to the magnetic property is saturated. Further, cooling may be started immediately after the temperature reaches the holding temperature, or cooling may also be started after the temperature is held for 6000 minutes or shorter. When this condition is satisfied, the achievement of high accumulation of the seeds oriented in the {200} plane further progresses to make it possible to more securely bring the accumulation degree of the {200} planes in the α-Fe phase to 30% or more after the cooling.

[Cooling after the Heating and Diffusion Treatment]

After the diffusion treatment, when the cooling is performed while the region that is not alloyed with the ferrite-forming element is remaining, as shown in FIG. 4B, at the time of the transformation to the α phase from the γ phase, the unalloyed region is transformed while taking over the crystal orientation of the regions in which the crystal grains 9 in the α phase oriented in {100} are already formed. Thereby, the accumulation degree of the {200} planes increases, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {222} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained, the crystal satisfying the condition of the Z value grows, and a high magnetic flux density can be obtained in an arbitrary direction in the metal sheet plane.

Further, as shown in FIG. 4C, when the A3 point or higher is held until the whole metal sheet is alloyed and the grain structures oriented in {100} reach the center of the metal sheet, as shown in FIG. 4D, the metal sheet is cooled as it is, and the texture in which the grain structures oriented in {100} reach the center of the metal sheet can be obtained. Thereby, the whole metal sheet is alloyed with the different metal, and the metal sheet having the texture in which the accumulation degree of the {200} planes in the α-Fe phase is not less than 30% nor more than 99% and the accumulation degree of the {222} planes in the α-Fe phase is not less than 0.01% nor more than 30% is obtained.

As above, the value of the accumulation degree of the {200} planes and the remaining state of the different metal on the surface of the base metal sheet change depending on the holding time of the temperature of the A3 point or higher and the holding temperature. The example shown in FIG. 4B is in a state where the grain structures oriented in {100} do not reach up to the center of the metal sheet and the different metal also remains on the surfaces, but it is also possible to obtain the grain structures oriented in {100} up to the center of the metal sheet and to alloy all the second layers on the surfaces.

Incidentally, at the time of the cooling after the diffusion treatment, a cooling rate is preferably not less than 0.1° C./sec nor more than 500° C./sec. When the cooling rate is less than 0.1° C./sec, a long time is required for the cooling, which is not appropriate, and when the cooling rate is greater than 500° C./sec, the metal sheet is sometimes deformed, and thus the cooling rate is preferably 500° C./sec or less.

Incidentally, when the second layers are made to remain on the obtainable Fe-based metal sheet having a thickness of not less than 10 μm nor more than 6 mm, the thickness of the second layer is preferably set to not less than 0.01 μm nor more than 500 μm. Further, a ratio of the α single phase region alloyed at this stage is preferably 1% or more in a cross section of the Fe-based metal sheet.

Further, it is also possible to form a structure as shown in FIG. 5, and in this case, an average cooling rate is set to satisfy the condition similar to that of the first embodiment, and thereby the above can be achieved.

EXAMPLE

Next, there will be explained experiments conducted by the present inventors. Conditions and the like in these experiments are examples employed for confirming the applicability and effects of the present invention, and the present invention is not limited to these examples.

Example 1

In this example, base metal sheets of No. 1 to No. 17 each made of a component A or B shown in Table 1 below were manufactured under various rolling conditions, to then have various different metals applied thereto as a second layer, and then Fe-based metal sheets were fabricated, of which the previously described Z value (=(A+0.97B)/0.98C) and the magnetic flux density difference ΔB were examined. Further, the relationship between various manufacturing conditions and an accumulation degree of {200} planes was also examined. Further, effects obtained by changing a starting temperature in an α-region rolling process were also examined in detail.

TABLE 1 COMPONENT A3 ELEMENT MASS % SERIES POINT C Si Mn Al P N S O OTHER A 925 0.0008 0.3 0.3 0.5 0.0003 0.0002 <0.0004 0.0002 — B 1010 0.0012 1.1 0.8 0.1 0.0002 0.0003 <0.0004 0.0001 — C 915 0.0032 0.2 0.08 0.05 0.0001 0.0003 <0.0004 0.0001 — D 870 0.0041 0.1 1.5 0.2 0.0001 0.0002 <0.0004 0.0001 — E 942 0.0105 0.2 0.5 0.7 0.0001 0.0003 <0.0004 0.0001 Cr: 0.5

First, ingots each having the component A or B shown in Table 1 and a balance being composed of Fe and inevitable impurities were melted by vacuum melting. Then, these were used as rolling materials to be worked into cold-rolled sheets (the base metal sheets) each having a predetermined thickness under conditions of hot rolling, α-region rolling, and cold rolling shown in Table 2 below.

TABLE 2 α-REGION HOT ROLLING ROLLING START FINISH START FINISH BASE TEMPER- THICK- TEMPER- THICK- REDUC- TEMPER- THICK- TEMPER- MATERIAL A3 ATURE NESS ATURE NESS TION ATURE NESS ATURE No. COMPONENT POINT ° C. mm ° C. mm RATIO ° C. mm ° C. 1 A 925 1150 250 1000 10 −3.22 950 10 920 2 A 925 1150 250 1000 10 −3.22 920 10 830 3 A 925 1150 250 1000 10 −3.22 850 10 830 4 A 925 1150 250 1000 10 −3.22 750 10 730 5 A 925 1150 250 1000 10 −3.22 650 10 640 6 A 925 1150 250 1000 10 −3.22 550 10 540 7 A 925 1150 250 1000 10 −3.22 450 10 450 8 A 925 1150 250 1000 10 −3.22 300 10 350 9 A 925 1150 250 1000 10 −3.22 250 10 250 10 B 1010 1200 280 1050 50 −1.72 1050 50 980 11 B 1010 1200 280 1050 50 −1.72 950 50 880 12 B 1010 1200 280 1050 50 −1.72 850 50 770 13 B 1010 1200 280 1050 50 −1.72 750 50 660 14 B 1010 1200 280 1050 50 −1.72 600 50 580 15 B 1010 1200 280 1050 50 −1.72 450 50 485 16 B 1010 1200 280 1050 50 −1.72 300 50 390 17 B 1010 1200 280 1050 50 −1.72 250 50 230 α-REGION ROLLING COLD ROLLING FINISH START FINISH TOTAL REDUCTION BASE THICK- REDUC- THICK- THICK- REDUC- REDUC- RATIO OF α MATERIAL NESS TION NESS NESS TION TION REGION + No. mm RATIO mm mm RATIO RATIO COLD ROLLING 1 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 2 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 3 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 4 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 5 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 6 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 7 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 8 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 9 2.5 −1.39 2.5 0.2 −2.53 −7.13 −3.91 10 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 11 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 12 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 13 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 14 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 15 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 16 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61 17 3.0 −2.81 3.0 0.5 −1.79 −6.33 −4.61

In the case of the component A, the ingots each having a thickness of 250 mm heated to 1150° C. were first subjected to hot rolling at a reduction ratio of −3.22 in terms of true strain, and hot-rolled sheets each having a thickness of 10 mm were obtained. Next, these hot-rolled sheets were each subjected to α-region rolling at a reduction ratio of −1.39 in terms of true strain at a temperature of 300 to 1000° C. These rolled sheets obtained by the α-region rolling were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −2.53 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.2 mm.

In the case of the component B, the ingots each having a thickness of 280 mm heated to 1200° C. were first subjected to hot rolling at a reduction ratio of −1.72 in terms of true strain, and hot-rolled sheets each having a thickness of 50 mm were obtained. Next, these hot-rolled sheets were each subjected to α-region rolling at a reduction ratio of −2.81 in terms of true strain at a temperature of 300 to 1050° C. These rolled sheets obtained by the α-region rolling were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −1.79 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.5 mm.

With respect to the base metal sheets obtained by the above procedure, a texture of a surface layer portion of each of the base materials was measured by X-ray diffraction to obtain an accumulation degree of {200} planes and an accumulation degree of {222} planes by the previously described method. Further, thinning was performed so that a structure could be observed from a direction perpendicular to an L cross section, and a region up to ¼t (t represents a thickness) from the surface was observed. The main phase of each of the obtained base metal sheets at room temperature was an α-Fe phase. Further, as a result of measurement, the A3 point at which the α-γ transformation occurred was 925° C. in the component A and 1010° C. in the component B.

Next, both surfaces of each of the base metal sheets of No. 1 to No. 17 shown in Table 2 were coated with each of various different metal elements as the second layer by a vapor deposition method, a sputtering method, or an electroplating method. As shown in Table 3 and Table 4 below, as the different metal element, any one of Al, Si, Mo, Ga, Sn, Ti, Ge, Sb, V, and W was selected. The thickness of each of the coatings was as shown in Table 3 and Table 4.

Next, an experiment was performed in which a heat treatment was performed on the base metal sheets to each of which the second layers were bonded under various conditions. A gold image furnace was used for the heat treatment, and a holding time was controlled by program control. During which the temperature increased to be held, the heat treatment was performed in an atmosphere vacuumed to a pressure of 10⁻³ Pa level. At the time of cooling, in the case of a cooling rate of 1° C./sec or lower, temperature control was performed in a vacuum by furnace output control. Further, in the case of the cooling rate of 10° C./sec or more, an Ar gas was introduced and the cooling rate was controlled by adjustment of its flow rate.

Here, there was examined a change in the texture among a temperature increasing process of heating up to the A3 point, a holding process of heating to a temperature of not lower than the A3 point nor higher than 1300° C. and holding the temperature, and a cooling process of cooling to a temperature of lower than the A3 point. Specifically, three base metal sheets with the same combination of the base material-coating conditions were prepared, of which a change in the texture was examined by performing a heat treatment experiment in each of the processes.

A sample for the temperature increasing process was fabricated in such a manner that the base metal sheet was heated from room temperature to the A3 point at a predetermined temperature increasing rate and was cooled to room temperature without any holding time. The cooling rate was set to 100° C./sec. The texture was measured by the method using the previously described X-ray diffraction method, and the X-ray was emitted from its surface, and the accumulation degree of {200} planes in the α-Fe phase and the accumulation degree of {222} planes in the α-Fe phase were obtained in an inverse pole figure.

A sample for the holding process was fabricated in such a manner that the base metal sheet was heated from room temperature to a predetermined temperature over the A3 point at a predetermined temperature increasing rate and was cooled to room temperature after a predetermined holding time. Then, the texture of the fabricated sample was measured in the same manner, and the accumulation degrees of {200} and {222} planes in the α-Fe phase were obtained.

A sample for the cooling process was fabricated in such a manner that the base metal sheet was heated from room temperature to a predetermined temperature over the A3 point at a predetermined temperature increasing rate and was cooled to room temperature at a predetermined cooling rate after a predetermined holding time. Further, in order to evaluate the accumulation degrees of {200} and {222} planes at an unalloyed position, a test piece was fabricated by removing a layer from the surface of the fabricated sample to a predetermined distance so that the unalloyed position might become an evaluation surface. Incidentally, when the whole metal sheet was alloyed, the evaluation surface was set to a position of ½ of the sheet thickness. With regard to the measurement of the texture of the fabricated sample, the X-ray was emitted from the surface of the test piece and from a predetermined surface of the test piece from which the layer was removed, and the accumulation degrees of {200} and {222} planes in the α-Fe phase of the surfaces were obtained in the same manner.

Next, magnetometry was performed in order to evaluate obtained products. First, the average magnetic flux density B50 to a magnetizing force of 5000 A/m and the magnetic flux density difference ΔB were obtained by using a SST (Single Sheet Tester). At this time, a measurement frequency was set to 50 Hz. When the average magnetic flux density B50 was obtained, as shown in FIG. 1, the magnetic flux density B50 was obtained every 22.5° in a circumferential direction of the product and an average value of the magnetic flux densities B50 in these 16 directions was calculated. Further, of the magnetic flux densities B50 in these 16 directions, the difference between the maximum value and the minimum value was set to the magnetic flux density difference ΔB. Next, the saturation magnetic flux density Bs was obtained by using a VSM (Vibrating Sample Magnetometer). The magnetizing force applied at this time was 0.8×10⁶ A/m. An evaluation value was set to the ratio B50/Bs of the average magnetic flux density B50 to the saturation magnetic flux density.

Further, by the previously described X-ray diffraction, intensity ratios of {001}<470>, {116}<6 12 1>, and {223}<692> were calculated, and thereby the previously described Z value was calculated.

Table 3 and Table 4 below show the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured in the respective processes during the manufacture and after the manufacture, the Z values of the obtained Fe-based metal sheets, and evaluation results of the magnetometry.

TABLE 3 MANUFACTURE SEEDING BASE MATERIAL SECOND LAYER TEMPERATURE SEEDING SEEDING INCREASING MEASURED ACCUMULATION ACCUMULATION THICKNESS THICKNESS BONDING RATE TEMPER- DEGREE OF {200} DEGREE OF{222} No. No. mm ELEMENT μm METHOD ° C./SEC ATURE PLANES IN αFe PLANES IN αFe  1 1 0.2 Al 4 VAPOR 20 925 14 13 DEPOSITION  2 2 0.2 Al 4 VAPOR 20 925 19 12 DEPOSITION  3 3 0.2 Al 4 VAPOR 20 925 25 10 DEPOSITION  4 4 0.2 Al 4 VAPOR 20 925 27 9.6 DEPOSITION  5 5 0.2 Al 4 VAPOR 20 925 33 9.1 DEPOSITION  6 6 0.2 Al 4 VAPOR 20 925 34 8.8 DEPOSITION  7 7 0.2 Al 4 VAPOR 20 925 35 8.7 DEPOSITION  8 8 0.2 Al 4 VAPOR 20 925 28 9.5 DEPOSITION  9 9 0.2 Al 4 VAPOR 20 925 27 9.8 DEPOSITION 10 1 0.2 Si 5 SPUTTERING 70 925 14 13 11 2 0.2 Si 5 SPUTTERING 70 925 18 12 12 3 0.2 Si 5 SPUTTERING 70 925 26 10 13 4 0.2 Si 5 SPUTTERING 70 925 27 9.4 14 5 0.2 Si 5 SPUTTERING 70 925 32 8.8 15 6 0.2 Si 5 SPUTTERING 70 925 35 8.2 16 7 0.2 Si 5 SPUTTERING 70 925 35 8.3 17 8 0.2 Si 5 SPUTTERING 70 925 29 9.7 18 9 0.2 Si 5 SPUTTERING 70 925 28 10 19 1 0.2 Mb 1.5 SPUTTERING 10 925 15 12 20 2 0.2 Mb 1.5 SPUTTERING 10 925 17 11 21 3 0.2 Mb 1.5 SPUTTERING 10 925 26 9.8 22 4 0.2 Mb 1.5 SPUTTERING 10 925 28 9.2 23 5 0.2 Mb 1.5 SPUTTERING 10 925 33 8.6 24 6 0.2 Mb 1.5 SPUTTERING 10 925 35 8.3 25 7 0.2 Mb 1.5 SPUTTERING 10 925 35 8.1 26 8 0.2 Mb 1.5 SPUTTERING 10 925 27 10 27 9 0.2 Mb 1.5 SPUTTERING 10 925 26 11 28 1 0.2 Ga 3 VAPOR 0.5 925 15 13 DEPOSITION 29 2 0.2 Ga 3 VAPOR 0.5 925 17 12 DEPOSITION 30 3 0.2 Ga 3 VAPOR 0.5 925 26 10 DEPOSITION 31 4 0.2 Ga 3 VAPOR 0.5 925 28 9.3 DEPOSITION 32 5 0.2 Ga 3 VAPOR 0.5 925 34 8.7 DEPOSITION 33 6 0.2 Ga 3 VAPOR 0.5 925 35 8.1 DEPOSITION 34 7 0.2 Ga 3 VAPOR 0.5 925 35 7.8 DEPOSITION 35 8 0.2 Ga 3 VAPOR 0.5 925 27 10 DEPOSITION 36 9 0.2 Ga 3 VAPOR 0.5 925 25 11 DEPOSITION 37 1 0.2 Sn 6 ELECTROLYTIC 5 925 16 13 PLATING 38 2 0.2 Sn 6 ELECTROLYTIC 5 925 19 11 PLATING 39 3 0.2 Sn 6 ELECTROLYTIC 5 925 27 9.5 PLATING 40 4 0.2 Sn 6 ELECTROLYTIC 5 925 28 9.1 PLATING 41 5 0.2 Sn 6 ELECTROLYTIC 5 925 32 8.7 PLATING 42 6 0.2 Sn 6 ELECTROLYTIC 5 925 33 8.3 PLATING 43 7 0.2 Sn 6 ELECTROLYTIC 5 925 34 8.2 PLATING 44 8 0.2 Sn 6 ELECTROLYTIC 5 925 27 9.7 PLATING 45 9 0.2 Sn 6 ELECTROLYTIC 5 925 26 11 PLATING SAVING AND ACHIEVEMENT OF HIGH ACCUMULATION ACHIEVEMENT OF HIGH ACHIEVEMENT OF HIGH GROWTH ACCUMULATION ACCUMULATION 1/2t 1/2t HOLDING ACCUMULATION ACCUMULATION ACCUMULATION ACCUMULATION TEMPERATURE HOLDING DEGREE OF {200} DEGREE OF {222} COOLING RATE DEGREE DEGREE No. ° C. TIME sec PLANES IN αFe PLANES IN αFe ° C./SEC OF {200} PLANES OF {222} PLANES  1 1000 20 16 13 150 16 13  2 1000 20 25 10.4 150 25 10.4  3 1000 20 30 9.1 150 30 9.1  4 1000 20 41 3.4 150 41 3.4  5 1000 20 53 1.8 150 53 1.8  6 1000 20 52 2.1 150 52 2.1  7 1000 20 50 2.3 150 50 2.3  8 1000 20 38 3.8 150 38 3.8  9 1000 20 37 4.2 150 37 4.2 10 1050 10 17 12 250 17 12 11 1050 10 24 12 250 24 12 12 1050 10 31 8.7 250 31 8.7 13 1050 10 42 3.2 250 42 3.2 14 1050 10 55 1.7 250 55 1.7 15 1050 10 54 1.9 250 54 1.9 16 1050 10 51 2.2 250 51 2.2 17 1050 10 39 4.1 250 39 4.1 18 1050 10 37 4.5 250 37 4.5 19 1250 10 15 13 10 15 13 20 1250 10 23 13 10 23 13 21 1250 10 30 9.3 10 30 9.3 22 1250 10 41 4.1 10 41 4.1 23 1250 10 52 2.4 10 52 2.4 24 1250 10 52 2.6 10 52 2.6 25 1250 10 51 2.9 10 51 2.9 26 1250 10 38 4.8 10 38 4.8 27 1250 10 37 5.5 10 37 5.5 28 980 100 17 11 50 17 11 29 980 100 27 9.8 50 27 9.8 30 980 100 33 8.5 50 33 8.5 31 980 100 43 3.5 50 43 3.5 32 980 100 57 1.8 50 57 1.8 33 980 100 56 2.1 50 56 2.1 34 980 100 55 2.3 50 55 2.3 35 980 100 40 3.9 50 40 3.9 36 980 100 37 4.5 50 37 4.5 37 1100 20 16 13 350 16 13 38 1100 20 27 11 350 27 11 39 1100 20 32 9.4 350 32 9.4 40 1100 20 45 3.1 350 45 3.1 41 1100 20 58 1.4 350 58 1.4 42 1100 20 57 1.9 350 57 1.9 43 1100 20 56 2.1 350 56 2.1 44 1100 20 43 3.8 350 43 3.8 45 1100 20 41 5.1 350 41 5.1 PRODUCT TEXTURE EVALUATION ACCUMULATION ACCUMULATION MAGNETIC FLUX DENSITY EVALUATION DEGREE OF DEGREE OF B50 Bs ΔB α-REGION No. {200} PLANES {222} PLANES Z T T B50/Bs T NOTE TEMPERATURE 1 16 13 1.2 1.60 2.05 0.78 0.070 COMPARITVE 950 EXAMPLE 1 2 25 10.4 2.1 1.66 2.05 0.81 0.065 PRESENT INVENTION 920 EXAMPLE 1 3 30 9.1 5.8 1.71 2.05 0.83 0.060 PRESENT INVENTION 850 EXAMPLE 2 4 41 3.4 23 1.77 2.05 0.86 0.056 PRESENT 750 EXAMPLE 3 5 53 1.8 160 1.84 2.05 0.90 0.018 PRESENT 650 EXAMPLE 4 6 52 2.1 120 1.87 2.05 0.91 0.021 PRESENT INVENTION 550 EXAMPLE 5 7 50 2.3 42 1.86 2.05 0.91 0.070 PRESENT INVENTION 450 EXAMPLE 6 8 38 3.8 3.5 1.80 2.05 0.88 0.145 PRESENT INVENTION 300 EXAMPLE 7 9 37 4.2 1.1 1.78 2.05 0.87 0.220 COMPARITVE 250 EXAMPLE 2 10 17 12 1.4 1.60 2.05 0.78 0.080 COMPARITVE 950 EXAMPLE 3 11 24 12 2.5 1.65 2.05 0.80 0.074 PRESENT INVENTION 920 EXAMPLE 8 12 31 8.7 3.8 1.66 2.05 0.81 0.070 PRESENT INVENTION 850 EXAMPLE 9 13 42 3.2 27 1.79 2.05 0.87 0.054 PRESENT INVENTION 750 EXAMPLE 10 14 55 1.7 156 1.88 2.05 0.92 0.015 PRESENT INVENTION 650 EXAMPLE 11 15 54 1.9 134 1.87 2.05 0.91 0.025 PRESENT INVENTION 550 EXAMPLE 12 16 51 2.2 51 1.86 2.05 0.91 0.086 PRESENT INVENTION 450 EXAMPLE 13 17 39 4.1 5.8 1.81 2.05 0.88 0.145 PRESENT INVENTION 300 EXAMPLE 14 18 37 4.5 1.7 1.74 2.05 0.85 0.210 COMPARITVE 250 EXAMPLE 4 19 15 13 1.3 1.59 2.05 0.78 0.087 COMPARITVE 950 EXAMPLE 5 20 23 13 2.4 1.66 2.05 0.81 0.081 PRESENT INVENTION 920 EXAMPLE 15 21 30 9.3 5.8 1.72 2.05 0.84 0.080 PRESENT INVENTION 850 EXAMPLE 16 22 41 4.1 19 1.78 2.05 0.87 0.074 PRESENT INVENTION 750 EXAMPLE 17 23 52 2.4 149 1.86 2.05 0.91 0.021 PRESENT INVENTION 650 EXAMPLE 18 24 52 2.6 174 1.86 2.05 0.91 0.018 PRESENT INVENTION 550 EXAMPLE 19 25 51 2.9 39 1.86 2.05 0.91 0.093 PRESENT INVENTION 450 EXAMPLE 20 26 38 4.8 3.1 1.77 2.05 0.86 0.138 PRESENT INVENTION 300 EXAMPLE 21 27 37 5.5 1.1 1.76 2.05 0.86 0.190 COMPARITVE 250 EXAMPLE 6 28 17 11 1.2 1.61 2.05 0.79 0.082 COMPARITVE 950 EXAMPLE 7 29 27 9.8 2.5 1.65 2.05 0.80 0.073 PRESENT INVENTION 920 EXAMPLE 22 30 33 8.5 8.5 1.73 2.05 0.84 0.073 PRESENT INVENTION 850 EXAMPLE 23 31 43 3.5 34 1.78 2.05 0.87 0.064 PRESENT INVENTION 750 EXAMPLE 24 32 57 1.8 112 1.87 2.05 0.91 0.017 PRESENT INVENTION 650 EXAMPLE 25 33 56 2.1 110 1.88 2.05 0.92 0.013 PRESENT INVENTION 550 EXAMPLE 26 34 55 2.3 74 1.87 2.05 0.91 0.087 PRESENT INVENTION 450 EXAMPLE 27 35 40 3.9 2.1 1.76 2.05 0.86 0.139 PRESENT INVENTION 300 EXAMPLE 28 36 37 4.5 0.6 1.74 2.05 0.85 0.210 COMPARITVE 250 EXAMPLE 8 37 16 13 0.8 1.60 2.05 0.78 0.086 COMPARITVE 950 EXAMPLE 9 38 27 11 2.2 1.65 2.05 0.80 0.079 PRESENT INVENTION 920 EXAMPLE 29 39 32 9.4 8.2 1.73 2.05 0.84 0.079 PRESENT INVENTION 850 EXAMPLE 30 40 45 3.1 34 1.81 2.05 0.88 0.065 PRESENT INVENTION 750 EXAMPLE 31 41 58 1.4 158 1.87 2.05 0.91 0.013 PRESENT INVENTION 650 EXAMPLE 32 42 57 1.9 189 1.89 2.05 0.92 0.009 PRESENT INVENTION 550 EXAMPLE 33 43 56 2.1 48 1.88 2.05 0.92 0.091 PRESENT INVENTION 450 EXAMPLE 34 44 43 3.8 2.7 1.77 2.05 0.86 0.136 PRESENT INVENTION 300 EXAMPLE 35 45 41 5.1 1.4 1.76 2.05 0.86 0.192 COMPARITVE 250 EXAMPLE 9

TABLE 4 MANUFACTURE SEEDING SEEDING SEEDING ACCUMULATION ACCUMULATION BASE MATERIAL SECOND LAYER TEMPERATURE DEGREE DEGREE THICKNESS THICKNESS INCREASING RATE MEASURED OF {200} PLANES OF {222} PLANES No. No. mm ELEMENT μm BONDING METHOD ° C./SEC TEMPERATURE IN αFe IN αFe 46 10 0.5 Ti 10 SPUTTERING 50 1010 14 14 47 11 0.5 Ti 10 SPUTTERING 50 1010 25 8.6 48 12 0.5 Ti 10 SPUTTERING 50 1010 30 1.9 49 13 0.5 Ti 10 SPUTTERING 50 1010 36 0.7 50 14 0.5 Ti 10 SPUTTERING 50 1010 38 0.3 51 15 0.5 Ti 10 SPUTTERING 50 1010 38 0.8 52 16 0.5 Ti 10 SPUTTERING 50 1010 27 3.8 53 17 0.5 Ti 10 SPUTTERING 50 1010 26 4.9 54 10 0.5 Ge 12 SPUTTERING 100 1010 13 13 55 11 0.5 Ge 12 SPUTTERING 100 1010 26 8.3 56 12 0.5 Ge 12 SPUTTERING 100 1010 30 1.9 57 13 0.5 Ge 12 SPUTTERING 100 1010 34 0.8 58 14 0.5 Ge 12 SPUTTERING 100 1010 37 0.3 59 15 0.5 Ge 12 SPUTTERING 100 1010 37 0.7 60 16 0.5 Ge 12 SPUTTERING 100 1010 26 3.9 61 17 0.5 Ge 12 SPUTTERING 100 1010 25 4.5 62 10 0.5 Sb 15 SPUTTERING 1 1010 12 12 63 11 0.5 Sb 15 SPUTTERING 1 1010 26 9.2 64 12 0.5 Sb 15 SPUTTERING 1 1010 30 2.4 65 13 0.5 Sb 15 SPUTTERING 1 1010 35 0.9 66 14 0.5 Sb 15 SPUTTERING 1 1010 37 0.4 67 15 0.5 Sb 15 SPUTTERING 1 1010 37 1.1 68 16 0.5 Sb 15 SPUTTERING 1 1010 26 4.1 69 17 0.5 Sb 15 SPUTTERING 1 1010 25 5.2 70 10 0.5 V 18 SPUTTERING 300 1010 13 13 71 11 0.5 V 18 SPUTTERING 300 1010 26 10 72 12 0.5 V 18 SPUTTERING 300 1010 31 8.5 73 13 0.5 V 18 SPUTTERING 300 1010 36 7.7 74 14 0.5 V 18 SPUTTERING 300 1010 38 6.3 75 15 0.5 V 18 SPUTTERING 300 1010 37 6.7 76 16 0.5 V 18 SPUTTERING 300 1010 27 10 77 17 0.5 V 18 SPUTTERING 300 1010 26 11 78 10 0.5 W 10 SPUTTERING 50 1010 13 14 79 11 0.5 W 10 SPUTTERING 50 1010 26 11 80 12 0.5 W 10 SPUTTERING 50 1010 31 9.1 81 13 0.5 W 10 SPUTTERING 50 1010 35 7.9 82 14 0.5 W 10 SPUTTERING 50 1010 37 6.4 83 15 0.5 W 10 SPUTTERING 50 1010 38 6.4 84 16 0.5 W 10 SPUTTERING 50 1010 26 10 85 17 0.5 W 10 SPUTTERING 50 1010 25 11 SAVING AND ACHIEVEMENT OF HIGH ACCUMULATION ACHIEVEMENT ACHIEVEMENT OF HIGH OF HIGH GROWTH ACCUMULATION ACCUMULATION ACCUMULATION ACCUMULATION 1/2t 1/2t HOLDING DEGREE OF DEGREE OF COOLING ACCUMULATION ACCUMULATION TEMPERATURE HOLDING {200} PLANES {222} PLANES RATE DEGREE OF DEGREE OF No. ° C. TIME sec IN αFe IN αFe ° C./SEC {200} PLANES {222} PLANES 46 1100 10 15 14 50 15 14 47 1100 10 32 8.6 50 32 8.6 48 1100 10 52 1.9 50 52 1.9 49 1100 10 67 0.7 50 67 0.7 50 1100 10 71 0.3 50 71 0.3 51 1100 10 68 0.8 50 68 0.8 52 1100 10 44 3.8 50 44 3.8 53 1100 10 40 4.9 50 40 4.9 54 1250 30 16 13 150 16 13 55 1250 30 33 8.3 150 33 8.3 56 1250 30 51 1.9 150 51 1.9 57 1250 30 65 0.8 150 65 0.8 58 1250 30 70 0.3 150 70 0.3 59 1250 30 67 0.7 150 67 0.7 60 1250 30 43 3.9 150 43 3.9 61 1250 30 41 4.5 150 41 4.5 62 1050 100 16 12 20 16 12 63 1050 100 31 9.2 20 31 9.2 64 1050 100 50 2.4 20 50 2.4 65 1050 100 66 0.9 20 66 0.9 66 1050 100 69 0.4 20 69 0.4 67 1050 100 64 1.1 20 64 1.1 68 1050 100 42 4.1 20 42 4.1 69 1050 100 39 5.2 20 39 5.2 70 1150 200 14 14 5 14 14 71 1150 200 33 8.4 5 33 8.4 72 1150 200 53 1.5 5 53 1.5 73 1150 200 66 0.8 5 66 0.8 74 1150 200 70 0.4 5 70 0.4 75 1150 200 67 0.8 5 67 0.8 76 1150 200 44 3.4 5 44 3.4 77 1150 200 41 4.8 5 41 4.8 78 1300 500 17 12 250 17 12 79 1300 500 34 8.7 250 34 8.7 80 1300 500 54 1.4 250 54 1.4 81 1300 500 65 0.9 250 65 0.9 82 1300 500 68 0.6 250 68 0.6 83 1300 500 65 1 250 65 1 84 1300 500 43 3..5 250 43 3.5 85 1300 500 40 4.8 250 40 4.8 PRODUCT TEXTURE EVALUATION ACCUMULATION ACCUMULATION MAGNETIC FLUX DENSITY EVALUATION DEGREE OF DEGREE OF B50 Bs ΔB α-REGION No. {200} PLANES {222} PLANES Z T T B50/Bs T NOTE TEMPERATURE 46 15 14 0.5 1.59 2.02 0.79 0.090 COMPARITVE 1050 EXAMPLE 11 47 32 8.6 2.5 1.73 2.02 0.86 0.080 PRESENT INVENTION 950 EXAMPLE 36 48 52 1.9 12 1.78 2.02 0.88 0.080 PRESENT INVENTION 850 EXAMPLE 37 49 67 0.7 75 1.89 2.02 0.94 0.030 PRESENT INVENTION 750 EXAMPLE 38 50 71 0.3 143 1.93 2.02 0.96 0.015 PRESENT INVENTION 600 EXAMPLE 39 51 68 0.8 116 1.92 2.02 0.95 0.019 PRESENT INVENTION 450 EXAMPLE 40 52 44 3.8 2.9 1.76 2.02 0.87 0.115 PRESENT INVENTION 300 EXAMPLE 41 53 40 4.9 0.9 1.75 2.02 0.87 0.200 COMPARITVE 250 EXAMPLE 12 54 16 13 0.7 1.58 2.02 0.78 0.091 COMPARITVE 1050 EXAMPLE 13 55 33 8.3 2.9 1.72 2.02 0.85 0.083 PRESENT INVENTION 950 EXAMPLE 42 56 51 1.9 18 1.79 2.02 0.89 0.070 PRESENT INVENTION 850 EXAMPLE 43 57 65 0.8 83 1.87 2.02 0.93 0.070 PRESENT INVENTION 750 EXAMPLE 44 58 70 0.3 183 1.93 2.02 0.96 0.045 PRESENT INVENTION 600 EXAMPLE 45 59 67 0.7 127 1.93 2.02 0.96 0.031 PRESENT INVENTION 450 EXAMPLE 46 60 43 3.9 4.7 1.77 2.02 0.88 0.138 PRESENT INVENTION 300 EXAMPLE 47 61 41 4.5 1.2 1.75 2.02 0.87 0.190 COMPARITVE 250 EXAMPLE 14 62 16 12 1.1 1.59 2.02 0.79 0.090 COMPARITVE 1050 EXAMPLE 15 63 31 9.2 2.4 1.73 2.02 0.86 0.080 PRESENT INVENTION 950 EXAMPLE 48 64 50 2.4 15 1.78 2.02 0.88 0.080 PRESENT INVENTION 850 EXAMPLE 49 65 66 0.9 77 1.87 2.02 0.93 0.076 PRESENT INVENTION 750 EXAMPLE 50 66 69 0.4 125 1.92 2.02 0.95 0.060 PRESENT INVENTION 600 EXAMPLE 51 67 64 1.1 108 1.92 2.02 0.95 0.042 PRESENT INVENTION 450 EXAMPLE 52 68 42 4.1 2.6 1.77 2.02 0.88 0.138 PRESENT INVENTION 300 EXAMPLE 53 69 39 5.2 1.4 1.76 2.02 0.87 0.220 COMPARITVE 250 EXAMPLE 16 70 14 14 0.4 1.58 2.02 0.78 0.230 COMPARITVE 1050 EXAMPLE 17 71 33 8.4 2.9 1.72 2.02 0.85 0.135 PRESENT INVENTION 950 EXAMPLE 54 72 53 1.5 36 1.77 2.02 0.88 0.094 PRESENT INVENTION 850 EXAMPLE 55 73 66 0.8 98 1.87 2.02 0.93 0.075 PRESENT INVENTION 750 EXAMPLE 56 74 70 0.4 178 1.94 2.02 0.96 0.061 PRESENT INVENTION 600 EXAMPLE 57 75 67 0.8 47 1.94 2.02 0.96 0.042 PRESENT INVENTION 450 EXAMPLE 58 76 44 3.4 10.4 1.76 2.02 0.87 0.137 PRESENT INVENTION 300 EXAMPLE 59 77 41 4.8 1.2 1.75 2.02 0.87 0.230 COMPARITVE 250 EXAMPLE 18 78 17 12 0.9 1.59 2.02 0.79 0.210 COMPARITVE 1050 EXAMPLE 19 79 34 8.7 4.7 1.73 2.02 0.86 0.143 PRESENT INVENTION 950 EXAMPLE 60 80 54 1.4 45 1.78 2.02 0.88 0.090 PRESENT INVENTION 850 EXAMPLE 61 81 65 0.9 118 1.87 2.02 0.93 0.064 PRESENT INVENTION 750 EXAMPLE 62 82 68 0.6 159 1.92 2.02 0.95 0.020 PRESENT INVENTION 600 EXAMPLE 63 83 65 1 69 1.92 2.02 0.95 0.031 PRESENT INVENTION 450 EXAMPLE 64 84 43 3.5 3.7 1.76 2.02 0.87 0.120 PRESENT INVENTION 300 EXAMPLE 65 85 40 4.8 1.7 1.75 2.02 0.87 0.230 COMPARITVE 250 EXAMPLE 20

In each of present invention examples, it was possible to confirm that Z is not less than 2.0 nor more than 200, the magnetic flux density difference ΔB becomes a small value as compared to comparative examples, and a high magnetic flux density can be obtained thoroughly in an in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.80 or more is obtained.

Further, in the present invention examples, as shown in Table 2 to Table 4, it was possible to confirm that the {200} plane in the α-Fe phase is likely to be highly accumulated at each of the stages of the heat treatment.

Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and a ratio of the α single phase region to the L cross section is 1% or more.

When the Z value was not less than 2 nor more than 200 as defined in the present invention as above, it was possible to confirm that a high magnetic flux density is obtained thoroughly in the in-plane circumferential direction. Further, in order to obtain the Fe-based metal sheet as above, the α-region rolling was performed at a temperature of higher than 300° C. and lower than the A3 point between the hot rolling and the cold rolling, thereby making it possible to obtain an intended product.

In contrast to this, when the base metal sheets obtained by performing the rolling under the conditions not satisfying the requirements of the present invention were used, it was not possible to obtain a high magnetic flux density such as that in the present invention examples in the in-plane circumferential direction thoroughly.

Example 2

In this example, base metal sheets of No. 18 to No. 35 each made of a component C, D, or E shown in Table 1 were manufactured under various rolling conditions, to then have various different metals applied thereto as a second layer, and then Fe-based metal sheets were fabricated, of which the previously described Z value (=(A+0.97B)/0.98C) and the magnetic flux density difference ΔB were examined. Further, the relationship between various manufacturing conditions and an accumulation degree of {200} planes was also examined. Further, effects obtained by changing a starting temperature in an α-region rolling process were also examined in detail.

First, ingots each having the component C, D, or E shown in Table 1 and a balance being composed of Fe and inevitable impurities were melted by vacuum melting. Then, these were used as rolling materials to be worked into cold-rolled sheets (the base metal sheets) each having a predetermined thickness under conditions of hot rolling, α-region rolling, and cold rolling shown in Table 5 below.

TABLE 5 HOT ROLLING α-REGION ROLLING BASE START FINISH START MATERIAL A3 TEMPERATURE THICKNESS TEMPERATURE THICKNESS REDUCTION TEMPERATURE THICKNESS No. COMPONENT POINT °C. mm °C. mm RATIO °C. mm 18 C 915 1050 200 930 60 −1.20 700 60 19 C 915 1050 200 930 30 −1.90 700 30 20 C 915 1050 200 930 20 −220 700 20 21 C 915 1050 200 930 10 −3.00 700 10 22 C 915 1050 200 930 8 −3.22 700 8 23 C 915 1050 200 930 4 −3.91 700 4 24 D 870 1050 300 980 15 −8.00 650 15 25 D 1370 1050 150 980 15 −220 650 15 26 D 870 1050 75 930 15 −1.61 650 15 27 D 870 1050 50 930 15 −1.20 650 15 28 D 870 1050 20 930 15 −0.29 650 15 29 E 942 1200 240 1050 50 −1.57 750 30 30 E 942 1200 240 1050 50 −1.57 750 30 31 E 942 1200 240 1050 50 −1.57 750 30 32 E 942 1200 240 1050 50 −1.57 750 30 33 E 942 1200 240 1050 50 −1.57 750 30 34 E 942 1200 240 1050 50 −1.57 750 30 35 E 942 1200 240 1050 50 −1.57 750 30 α-REGION ROLLING COLD ROLLING BASE FINISH START FINISH TOTAL α REGION + MATERIAL TEMPERATURE THICKNESS REDUCTION THICKNESS THICKNESS REDUCTION REDUCTION COLD No. °C. mm RATIO mm mm RATIO RATIO ROLLING 18 610 2 −8.40 2 0.05 −1.74 −6.35 −5.14 19 610 2 −2.71 2 025 −1.74 −6.35 −4.45 20 610 2 −2.00 2 025 −1.74 −6.05 −4.05 21 610 2 −1.61 2 0.35 −1.74 −625 −3.35 22 610 2 −1.39 2 0.35 −1.74 −6.35 −0.13 23 610 2 −0.69 2 0.35 −1.74 −6.35 −2.44 24 570 3.5 −1.46 3.5 0.5 −1.95 −6.40 −3.40 25 570 8.5 −1.46 8.5 0.5 −1.95 −5.70 −8.40 26 570 3.5 −1.46 3.5 0.5 −1.95 −5.01 −8.40 27 570 3.5 −1.46 3.5 0.5 −1.95 −4.61 −3.40 28 570 3.5 −1.46 3.5 05 −1.95 −3.69 −3.40 29 670 6 −1.61 6 3 −0.69 −0.87 −2.30 30 670 6 −1.61 6 2 −1.10 −4.28 −2.71 31 670 6 −1.61 6 1 −1.79 −4.97 −3.40 32 670 6 −1.61 6 02 −3.40 −6.58 −5.01 33 670 6 −1.61 6 0.1 −4.09 −7.27 −5.70 34 670 6 −1.61 6 0.05 −4.79 −7.97 −6.40 35 670 6 −1.51 6 0.01 −6.40 −9.57 −8.01

In the case of the component C, first, the ingots each having a thickness of 200 mm heated to 1050° C. were each subjected to hot rolling at a reduction ratio of −1.20 to −3.91 in terms of true strain, and hot-rolled sheets each having a thickness of 4 mm to 60 mm were obtained. Next, α-region rolling was started at 700° C., and these hot-rolled sheets were each subjected to the α-region rolling at a reduction ratio of −0.69 to −3.40 in terms of true strain to a thickness of 2 mm. Then, these rolled sheets were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −1.74 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.35 mm.

In the case of the component D, first, the ingots each having a thickness of 20 mm to 300 mm heated to 1050° C. were each subjected to hot rolling at a reduction ratio of −0.29 to −3.00 in terms of true strain, and hot-rolled sheets each having a thickness of 15 mm were obtained. Next, α-region rolling was started at 650° C., and these hot-rolled sheets were each subjected to the α-region rolling at a reduction ratio of −1.46 in terms of true strain to a thickness of 3.5 mm. Then, these rolled sheets were pickled, and then the base metal sheets were obtained by cold rolling. At this time, the reduction ratio was −1.95 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.50 mm.

In the case of the component E, first, the ingots each having a thickness of 240 mm heated to 1200° C. were each subjected to hot rolling at a reduction ratio of −1.57 in terms of true strain, and hot-rolled sheets each having a thickness of 50 mm were obtained. Next, α-region rolling was started at 750° C., and these hot-rolled sheets were each subjected to the α-region rolling at a reduction ratio of −1.61 in terms of true strain to a thickness of 6.0 mm. Then, these rolled sheets were pickled, and then the base metal sheets were obtained by cold rolling. At this time, each of the reduction ratios was −0.69 to −6.40 in terms of true strain, and as a result, the thickness of each of the obtained base metal sheets was 0.01 mm to 3.0 mm.

With respect to the base metal sheets obtained by the above procedure, a texture of a surface layer portion of each of the base materials was measured by X-ray diffraction to obtain an accumulation degree of {200} planes and an accumulation degree of {222} planes by the previously described method. Further, thinning was performed so that a structure could be observed from a direction perpendicular to an L cross-section, and a region up to ¼t from the surface was observed. The main phase of each of the obtained base metal sheets at room temperature was an α-Fe phase. Further, as a result of measurement, the A3 point at which the α-γ transformation occurred was 915° C. in the component C, 870° C. in the component D, and 942° C. in the component E.

Next, both surfaces of each of the base metal sheets of No. 18 to No. 35 shown in Table 5 were coated with each of various different metal elements as the second layer by a vapor deposition method, a sputtering method, an electroplating method, or a hot dipping method. As shown in Table 6 and Table 7 below, as the different metal element, any one of Al, Si, Ga, Sn, V, W, Mo, and Zn was selected. The thickness of each of the coatings was as shown in Table 6 and Table 7.

Next, an experiment was performed in which a heat treatment was performed on the base metal sheets to each of which the second layers were bonded under various conditions. As a method of the experiment, the experiment was performed by the same method described in Example 1. Further, the observation of a texture in this period was also performed by the same method described in Example 1.

Further, magnetometry was performed in the same manner as that in Example 1 in order to evaluate obtained products, and further the Z value was calculated by X-ray diffraction.

Table 6 and Table 7 below show the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured in the respective processes during the manufacture and after the manufacture, the Z values of the obtained Fe-based metal sheets, and evaluation results of the magnetometry.

TABLE 6 MANUFACTURE SEEDING SEEDING SEEDING ACCUMULATION ACCUMULATION BASE MATERIAL SECOND LAYER TEMPERATURE DEGREE OF DEGREE OF THICKNESS THICKNESS INCREASING RATE MEASURED {200} PLANES {222} PLANES No. No. mm ELEMENT μm BONDING METHOD ° C./SEC TEMPERATURE IN αFe IN αFe 86 18 0.35 Al 7 VAPOR DEPOSITION 20 915 34 7.8 87 19 0.35 Al 7 VAPOR DEPOSITION 20 915 34 7.9 88 20 0.35 Al 7 VAPOR DEPOSITION 20 915 33 8.3 89 21 0.35 Al 7 VAPOR DEPOSITION 20 915 30 9.3 90 22 0.35 Al 7 VAPOR DEPOSITION 20 915 28 10 91 23 0.35 Al 7 VAPOR DEPOSITION 20 915 26 10 92 18 0.35 Si 8 SPUTTERING 10 915 35 7.1 93 19 0.35 Si 8 SPUTTERING 10 915 34 7.3 94 20 0.35 Si 8 SPUTTERING 10 915 32 7.8 95 21 0.35 Si 8 SPUTTERING 10 915 27 9.7 96 22 0.35 Si 8 SPUTTERING 10 915 26 10 97 23 0.35 Si 8 SPUTTERING 10 915 25 11 98 18 0.35 Ga 6 VAPOR DEPOSITION 0.5 915 32 8.5 99 19 0.35 Ga 6 VAPOR DEPOSITION 0.5 915 32 8.6 100 20 0.35 Ga 6 VAPOR DEPOSITION 0.5 915 31 8.9 101 21 0.35 Ga 6 VAPOR DEPOSITION 0.5 915 28 9.8 102 22 0.35 Ga 6 VAPOR DEPOSITION 0.5 915 26 10 103 23 0.35 Ga 6 VAPOR DEPOSITION 0.5 915 25 11 104 18 0.35 Sn 10 ELECTROLYTIC 5 915 36 6.5 PLATING 105 19 0.35 Sn 10 ELECTROLYTIC 5 915 35 6.7 PLATING 106 20 0.35 Sn 10 ELECTROLYTIC 5 915 36 7.2 PLATING 107 21 0.35 Sn 10 ELECTROLYTIC 5 915 31 8.7 PLATING 108 22 0.35 Sn 10 ELECTROLYTIC 5 915 27 9.3 PLATING 109 23 0.35 Sn 10 ELECTROLYTIC 5 915 25 11 PLATING 110 18 0.35 V 11 SPUTTERING 10 915 34 7.9 111 19 0.35 V 11 SPUTTERING 10 915 33 8.2 112 20 0.35 V 11 SPUTTERING 10 915 31 8.6 113 21 0.35 V 11 SPUTTERING 10 915 28 9.7 114 22 0.35 V 11 SPUTTERING 10 915 27 10 115 23 0.35 V 11 SPUTTERING 10 915 25 10 116 18 0.35 W 6 SPUTTERING 0.5 915 34 7.6 117 19 0.35 W 6 SPUTTERING 0.5 915 33 8.2 118 20 0.35 W 6 SPUTTERING 0.5 915 31 9.2 119 21 0.35 W 6 SPUTTERING 0.5 915 28 10 120 22 0.35 W 6 SPUTTERING 0.5 915 27 10 121 23 0.35 W 6 SPUTTERING 0.5 915 25 11 SAVING AND ACHIEVEMENT OF HIGH ACCUMULATION ACHIEVEMENT ACHIEVEMENT OF HIGH OF HIGH GROWTH ACCUMULATION ACCUMULATION ACCUMULATION ACCUMULATION DEGREE OF DEGREE OF 1/2t ACCUMULATION 1/2t ACCUMULATION HOLDING TEMPERATURE HOLDING {200} PLANES {222} PLANES COOLING RATE DEGREE OF DEGREE OF No. ° C. TIME sec IN αFe IN αFe ° C./SEC {200} PLANES {222} PLANES 86 1000 40 63 0.8 100 63 0.8 87 1000 40 62 0.8 100 62 0.8 88 1000 40 61 0.9 100 61 0.9 89 1000 40 50 3.9 100 47 3.9 90 1000 40 42 5.7 100 40 5.7 91 1000 40 32 8.5 100 32 8.5 92 1050 25 62 0.9 20 62 0.9 93 1050 25 61 0.9 20 61 0.9 94 1050 25 60 1.2 20 60 1.2 95 1050 25 45 3.5 20 45 3.5 96 1050 25 42 5.3 20 42 5.3 97 1050 25 31 9.1 20 31 9.1 98 950 120 60 1.1 50 60 1.1 99 950 120 60 1 50 60 1 100 950 120 59 1.2 50 59 1.2 101 950 120 43 4.5 50 43 4.5 102 950 120 41 6.2 50 41 6.2 103 950 120 30 9.7 50 30 9.7 104 1000 10 64 0.7 200 64 0.7 105 1000 10 64 0.7 200 64 0.7 106 1000 10 63 0.8 200 63 0.8 107 1000 10 50 2.8 200 50 2.8 108 1000 10 43 4.7 200 43 4.7 109 1000 10 33 7.9 200 33 7.9 110 1200 15 61 0.9 250 61 0.9 111 1200 15 61 0.9 250 61 0.9 112 1200 15 60 1.2 250 60 1.2 113 1200 15 45 4.3 250 45 4.3 114 1200 15 39 6.6 250 39 6.6 115 1200 15 30 8.2 250 30 8.2 116 1300 30 60 1.1 80 60 1.1 117 1300 30 59 1.3 80 59 1.3 118 1300 30 58 1.6 80 58 1.6 119 1300 30 46 3.5 80 46 3.5 120 1300 30 38 6.9 80 38 6.9 121 1300 30 31 9.1 80 31 9.1 PRODUCT TEXTURE EVALUATION ACCUMULATION ACCUMULATION MAGNETIC FLUX DENSITY EVALUATION DEGREE OF DEGREE OF B50 Bs ΔB α-REGION No. {200} PLANES {222} PLANES Z T T B50/Bs T NOTE TEMPERATURE 86 63 0.8 135 1.87 2.04 0.92 0.038 PRESENT INVENTION 700 EXAMPLE 66 87 62 0.8 120 1.86 2.04 0.91 0.048 PRESENT INVENTION 700 EXAMPLE 67 88 61 0.9 52 1.86 2.04 0.91 0.053 PRESENT INVENTION 700 EXAMPLE 68 89 47 3.9 24 1.82 2.04 0.89 0.068 PRESENT INVENTION 700 EXAMPLE 69 90 40 5.7 8.6 1.76 2.04 0.86 0.076 PRESENT INVENTION 700 EXAMPLE 70 91 32 8.5 2.5 1.72 2.04 0.84 0.091 PRESENT INVENTION 700 EXAMPLE 71 92 62 0.9 157 1.86 2.04 0.91 0.035 PRESENT INVENTION 700 EXAMPLE 72 93 61 0.9 132 1.86 2.04 0.91 0.046 PRESENT INVENTION 700 EXAMPLE 73 94 60 1.2 62 1.85 2.04 0.91 0.057 PRESENT INVENTION 700 EXAMPLE 74 95 45 3.5 28 1.78 2.04 0.87 0.066 PRESENT INVENTION 700 EXAMPLE 75 96 42 5.3 9.4 1.76 2.04 0.86 0.084 PRESENT INVENTION 700 EXAMPLE 76 97 31 9.1 3.8 1.73 2.04 0.85 0.091 PRESENT INVENTION 700 EXAMPLE 77 98 60 1.1 167 1.86 2.04 0.91 0.031 PRESENT INVENTION 700 EXAMPLE 78 99 60 1 121 1.86 2.04 0.91 0.047 PRESENT INVENTION 700 EXAMPLE 79 100 59 1.2 71 1.85 2.04 0.91 0.054 PRESENT INVENTION 700 EXAMPLE 80 101 43 4.5 31 1.79 2.04 0.88 0.068 PRESENT INVENTION 700 EXAMPLE 81 102 41 6.2 10 1.76 2.04 0.86 0.079 PRESENT INVENTION 700 EXAMPLE 82 103 30 9.7 2.6 1.72 2.04 0.84 0.087 PRESENT INVENTION 700 EXAMPLE 83 104 64 0.7 184 1.88 2.04 0.92 0.027 PRESENT INVENTION 700 EXAMPLE 84 105 64 0.7 137 1.87 2.04 0.92 0.044 PRESENT INVENTION 700 EXAMPLE 85 106 63 0.8 68 1.88 2.04 0.92 0.057 PRESENT INVENTION 700 EXAMPLE 86 107 50 2.8 32 1.81 2.04 0.89 0.071 PRESENT INVENTION 700 EXAMPLE 87 108 43 4.7 9.4 1.75 2.04 0.86 0.082 PRESENT INVENTION 700 EXAMPLE 88 109 33 7.9 3.1 1.71 2.04 0.84 0.094 PRESENT INVENTION 700 EXAMPLE 89 110 61 0.9 154 1.87 2.04 0.92 0.022 PRESENT INVENTION 700 EXAMPLE 90 111 61 0.9 118 1.87 2.04 0.92 0.039 PRESENT INVENTION 700 EXAMPLE 91 112 60 1.2 66 1.86 2.04 0.91 0.053 PRESENT INVENTION 700 EXAMPLE 92 113 45 4.3 24 1.81 2.04 0.89 0.067 PRESENT INVENTION 700 EXAMPLE 93 114 39 6.6 8.9 1.76 2.04 0.86 0.075 PRESENT INVENTION 700 EXAMPLE 94 115 30 8.2 4.2 1.72 2.04 0.84 0.088 PRESENT INVENTION 700 EXAMPLE 95 116 60 1.1 186 1.86 2.04 0.91 0.019 PRESENT INVENTION 700 EXAMPLE 96 117 59 1.3 136 1.85 2.04 0.91 0.032 PRESENT INVENTION 700 EXAMPLE 97 118 58 1.6 74 1.85 2.04 0.91 0.050 PRESENT INVENTION 700 EXAMPLE 98 119 46 3.5 28 1.79 2.04 0.88 0.062 PRESENT INVENTION 700 EXAMPLE 99 120 38 6.9 12 1.76 2.04 0.86 0.072 PRESENT INVENTION 700 EXAMPLE 100 121 31 9.1 3.9 1.71 2.04 0.84 0.093 PRESENT INVENTION 700 EXAMPLE 101

TABLE 7 MANUFACTURE SEEDING SEEDING SEEDING ACCUMULATION ACCUMULATION BASE MATERIAL SECOND LAYER TEMPERATURE DEGREE OF DEGREE OF THICKNESS THICKNESS INCREASING RATE MEASURED {200} PLANES {222} PLANES No. No. mm ELEMENT μm BONDING METHOD ° C./SEC TEMPERATURE IN αFe IN αFe 122 24 0.5 Al 10 VAPOR DEPOSITION 10 870 31 7.7 123 25 0.5 Al 10 VAPOR DEPOSITION 10 870 31 7.9 124 26 0.5 Al 10 VAPOR DEPOSITION 10 870 30 9.3 125 27 0.5 Al 10 VAPOR DEPOSITION 10 870 27 9.8 126 28 0.5 Al 10 VAPOR DEPOSITION 10 870 25 10 127 24 0.5 Si 12 VAPOR DEPOSITION 20 870 31 8.1 128 25 0.5 Si 12 VAPOR DEPOSITION 20 870 31 8.2 129 26 0.5 Si 12 VAPOR DEPOSITION 20 870 30 9.3 130 27 0.5 Si 12 VAPOR DEPOSITION 20 870 27 10 131 28 0.5 Si 12 VAPOR DEPOSITION 20 870 26 11 132 24 0.5 Mb 8 SPUTTERING 1 870 33 6.8 133 25 0.5 Mb 8 SPUTTERING 1 870 32 7.3 134 26 0.5 Mb 8 SPUTTERING 1 870 30 8.8 135 27 0.5 Mb 8 SPUTTERING 1 870 27 9.3 136 28 0.5 Mb 8 SPUTTERING 1 870 25 10 137 29 3 Al 120 HOT DIPPING 2 942 13 13 138 30 2 Al 80 HOT DIPPING 2 942 25 10 139 31 1 Al 40 HOT DIPPING 2 942 31 8.3 140 32 0.2 Al 8 VAPOR DEPOSITION 2 942 32 7.5 141 33 0.1 Al 4 VAPOR DEPOSITION 2 942 33 6.7 142 34 0.05 Al 2 VAPOR DEPOSITION 2 942 33 6.5 143 35 0.01 Al 0.4 VAPOR DEPOSITION 2 942 32 6.4 144 29 3 Sn 60 HOT DIPPING 5 942 12 12 145 30 2 Sn 40 HOT DIPPING 5 942 25 10 146 31 1 Sn 20 HOT DIPPING 5 942 32 8.1 147 32 0.2 Sn 4 ELECTROLYTIC 5 942 33 7.1 PLATING 148 33 0.1 Sn 2 ELECTROLYTIC 5 942 34 6.3 PLATING 149 34 0.05 Sn 1 ELECTROLYTIC 5 942 35 6.1 PLATING 150 35 0.01 Sn 0.2 ELECTROLYTIC 5 942 34 6.6 PLATING 151 29 3 Zn 60 HOT DIPPING 1 942 14 13 152 30 2 Zn 40 HOT DIPPING 1 942 25 11 153 31 1 Zn 20 HOT DIPPING 1 942 30 8.8 154 32 0.2 Zn 4 ELECTROLYTIC 1 942 31 7.8 155 33 0.1 Zn 2 ELECTROLYTIC 1 942 32 6.5 156 34 0.05 Zn 1 ELECTROLYTIC 1 942 32 6.3 157 35 0.01 Zn 0.4 ELECTROLYTIC 1 942 32 6.7 SAVING AND ACHIEVEMENT OF HIGH ACCUMULATION ACHIEVEMENT OF HIGH ACHIEVEMENT OF HIGH GROWTH ACCUMULATION ACCUMULATION ACCUMULATION ACCUMULATION 1/2t 1/2t DEGREE OF DEGREE OF ACCUMULATION ACCUMULATION HOLDING TEMPERATURE HOLDING {200} PLANES {222} PLANES COOLING RATE DEGREE OF DEGREE OF No. ° C. TIME sec IN αFe IN αFe ° C./SEC {200} PLANES {222} PLANES 122 930 20 56 1.6 80 56 1.6 123 930 20 55 1.8 80 55 1.8 124 930 20 52 2.5 80 52 2.5 125 930 20 40 5.9 80 40 5.9 126 930 20 32 9.3 80 32 9.3 127 980 60 54 1.7 20 54 1.7 128 980 60 53 1.9 20 53 1.9 129 980 60 51 2.8 20 51 2.8 130 980 60 39 7.4 20 39 7.4 131 980 60 33 9.5 20 33 9.5 132 1000 15 56 1.7 50 56 1.7 133 1000 15 56 1.7 50 56 1.7 134 1000 15 53 2.1 50 53 2.1 135 1000 15 41 6.3 50 41 6.3 136 1000 15 31 9.3 50 31 9.3 137 1050 25 15 13 100 15 13 138 1050 25 32 9.5 100 32 9.5 139 1050 25 50 2.8 100 50 2.8 140 1050 25 54 2.1 100 54 2.1 141 1050 25 55 1.8 100 55 1.8 142 1050 25 56 1.7 100 56 1.7 143 1050 25 55 1.8 100 55 1.8 144 1100 60 14 14 200 14 14 145 1100 60 32 9.4 200 32 9.4 146 1100 60 51 2.5 200 51 2.5 147 1100 60 56 1.8 200 56 1.8 148 1100 60 57 1.3 200 57 1.3 149 1100 60 57 1.1 200 57 1.1 150 1100 60 56 1.4 200 56 1.4 151 980 200 16 12 50 16 12 152 980 200 30 9.8 50 30 9.8 153 980 200 50 3.1 50 50 3.1 154 980 200 52 2.5 50 52 2.5 155 980 200 54 2.1 50 54 2.1 156 980 200 55 1.9 50 55 1.9 157 980 200 54 2.1 50 54 2.1 PRODUCT TEXTURE EVALUATION ACCUMULATION ACCUMULATION MAGNETIC FLUX DENSITY EVALUATION DEGREE OF DEGREE OF B50 Bs ΔB α-REGION No. {200} PLANES {222} PLANES Z T T B50/Bs T NOTE TEMPERATURE 122 56 1.6 98 1.85 1.98 0.93 0.024 PRESENT INVENTION 650 EXAMPLE 102 123 55 1.8 78 1.85 1.98 0.93 0.028 PRESENT INVENTION 650 EXAMPLE 103 124 52 2.5 57 1.83 1.98 0.92 0.041 PRESENT INVENTION 650 EXAMPLE 104 125 40 5.9 24 1.73 1.98 0.87 0.057 PRESENT INVENTION 650 EXAMPLE 105 126 32 9.3 3.8 1.69 1.98 0.85 0.087 PRESENT INVENTION 650 EXAMPLE 106 127 54 1.7 110 1.84 1.98 0.93 0.021 PRESENT INVENTION 650 EXAMPLE 107 128 53 1.9 76 1.84 1.98 0.93 0.025 PRESENT INVENTION 650 EXAMPLE 108 129 51 2.8 65 1.82 1.98 0.92 0.035 PRESENT INVENTION 650 EXAMPLE 109 130 39 7.4 29 1.76 1.98 0.89 0.053 PRESENT INVENTION 650 EXAMPLE 110 131 33 9.5 5.6 1.68 1.98 0.85 0.086 PRESENT INVENTION 650 EXAMPLE 111 132 56 1.7 105 1.86 1.98 0.94 0.023 PRESENT INVENTION 650 EXAMPLE 112 133 56 1.7 68 1.86 1.98 0.94 0.031 PRESENT INVENTION 650 EXAMPLE 113 134 53 2.1 59 1.84 1.98 0.93 0.045 PRESENT INVENTION 650 EXAMPLE 114 135 41 6.3 23 1.77 1.98 0.89 0.072 PRESENT INVENTION 650 EXAMPLE 115 136 31 9.3 4.2 1.68 1.98 0.85 0.092 PRESENT INVENTION 650 EXAMPLE 116 137 15 13 0.9 1.59 2.02 0.79 0.086 COMPARITIVE 750 EXAMPLE 21 138 32 9.5 2.5 1.73 2.02 0.86 0.062 PRESENT INVENTION 750 EXAMPLE 117 139 50 2.8 35 1.79 2.02 0.89 0.053 PRESENT INVENTION 750 EXAMPLE 118 140 54 2.1 65 1.83 2.02 0.91 0.041 PRESENT INVENTION 750 EXAMPLE 119 141 55 1.8 114 1.83 2.02 0.91 0.032 PRESENT INVENTION 750 EXAMPLE 120 142 56 1.7 126 1.83 2.02 0.91 0.018 PRESENT INVENTION 750 EXAMPLE 121 143 55 1.8 132 1.83 2.02 0.91 0.015 PRESENT INVENTION 750 EXAMPLE 122 144 14 14 0.4 1.60 2.02 0.79 0.092 COMPARITIVE 750 EXAMPLE 22 145 32 9.4 3.2 1.72 2.02 0.85 0.068 PRESENT INVENTION 750 EXAMPLE 123 146 51 2.5 29 1.76 2.02 0.87 0.052 PRESENT INVENTION 750 EXAMPLE 124 147 56 1.8 59 1.84 2.02 0.91 0.043 PRESENT INVENTION 750 EXAMPLE 125 148 57 1.3 94 1.84 2.02 0.91 0.029 PRESENT INVENTION 750 EXAMPLE 126 149 57 1.1 123 1.85 2.02 0.92 0.021 PRESENT INVENTION 750 EXAMPLE 127 150 56 1.4 135 1.84 2.02 0.91 0.018 PRESENT INVENTION 750 EXAMPLE 128 151 16 12 1.1 1.58 2.02 0.78 0.087 COMPARITIVE 750 EXAMPLE 23 152 30 9.8 4.5 1.71 2.02 0.85 0.058 PRESENT INVENTION 750 EXAMPLE 129 153 50 3.1 27 1.79 2.02 0.89 0.047 PRESENT INVENTION 750 EXAMPLE 130 154 52 2.5 49 1.83 2.02 0.91 0.039 PRESENT INVENTION 750 EXAMPLE 131 155 54 2.1 79 1.83 2.02 0.91 0.025 PRESENT INVENTION 750 EXAMPLE 132 156 55 1.9 132 1.83 2.02 0.91 0.018 PRESENT INVENTION 750 EXAMPLE 133 157 54 2.1 172 1.83 2.02 0.91 0.012 PRESENT INVENTION 750 EXAMPLE 134

In each of present invention examples, it was possible to confirm that the magnetic flux density difference ΔB becomes a small value as compared to comparative examples, and a high magnetic flux density is obtained thoroughly in an in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.86 or more is obtained.

Further, in the present invention examples, as shown in Table 5 to Table 7, it was possible to confirm that the {200} plane in the α-Fe phase is likely to be highly accumulated at each of the stages of the heat treatment.

Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and a ratio of the α single phase region to the L cross section is 1% or more.

When the Z value was not less than 2 nor more than 200 as defined in the present invention as above, it was possible to confirm that a high magnetic flux density is obtained thoroughly in the in-plane circumferential direction. Further, in order to obtain the Fe-based metal sheet as above, the α-region rolling was performed at a temperature of 300° C. or higher and lower than the A3 point between the hot rolling and the cold rolling, thereby making it possible to obtain an intended product.

In contrast to this, when the base metal sheets obtained by performing the α-region rolling under the condition not satisfying the requirements of the present invention were used, it was not possible to obtain a high magnetic flux density such as that in the present invention examples in the in-plane circumferential direction thoroughly.

Example 3

In this example, as base metal sheets, Fe-based metal sheets were fabricated in a manner that pure irons each containing C: 0.050 mass %, Si: 0.0001 mass %, and Al: 0.0002 mass %, and having a balance being composed of Fe and inevitable impurities were subjected to decarburization to have Al applied thereto as a second layer, of which the previously described Z value (=(A+0.97B)/0.98C) and the magnetic flux density difference ΔB were examined. Further, the relationship between manufacturing conditions and an accumulation degree of {200} planes was also examined.

First, ingots were melted by vacuum melting, and then were subjected to hot rolling and cold rolling to be worked to a predetermined thickness, and the base metal sheets each composed of the previously described composition were obtained. Incidentally, the A1 point of the base metal sheets was 727° C.

In the hot rolling, the ingots each having a thickness of 230 mm heated to 1000° C. were thinned down to a thickness of 50 mm, and hot-rolled sheets were obtained. Sheet materials having various thicknesses were cut out from these hot-rolled sheets by machining and then were subjected to the cold rolling, and thereby cold-rolled sheets each having a thickness of 8 μm to 750 μm (the base metal sheets) were obtained.

Incidentally, the main phase of each of the base metal sheets at room temperature was an α-Fe phase and as a result of measurement, the A3 point at which the α-γ transformation occurred was 911° C. Further, a texture in the α-Fe phase of each of the base metal sheets was measured by X-ray diffraction, and by the previously described method, an accumulation degree of {200} planes and an accumulation degree of {222} planes were obtained. Further, as a result that up to the cold rolling was performed, it was confirmed that of each of the base metal sheets, the accumulation degree of the {200} planes is 20 to 26% and the accumulation degree of the {222} planes is 18 to 24%.

Next, these base metal sheets were subjected to decarburization annealing so that a decarburized depth (a distance x) might become 1 μm to 59 μm. A decarburization condition was set that the temperature is 800° C. and the decarburization time is 0.05 minutes to 550 minutes. With regard to the atmosphere during the decarburization annealing, a strong decarburizing atmosphere was applied in the case of the decarburization annealing being performed for one minute or shorter, and in the case of the decarburization annealing being performed for longer than one minute, a weak decarburizing atmosphere was applied in the first half of the decarburization annealing and a strong decarburizing atmosphere was applied in the second half of the decarburization annealing.

Then, after the decarburization annealing was performed, the decarburized depth and the C content of a decarburized region were measured and a structure and a crystal orientation of a surface layer were examined. The measurement of the crystal orientation was performed by the method using the previously described X-ray diffraction method, the X-ray was emitted from the surface, and the accumulation degree of the {200} planes in the α-Fe phase was obtained.

After the decarburization annealing, both surfaces of each of the base metal sheets were coated with Al as the second layer by an ion plating method (hereinafter, an IP method) to each have a thickness of 1 μm.

Next, an experiment was performed in which a heat treatment was performed on the base metal sheets to each of which the second layers was bonded under various conditions. A gold image furnace was used for the heat treatment, and a temperature increasing rate, a holding temperature, and a holding time were variously controlled by program control. During which the temperature increased to be held, the heat treatment was performed in an atmosphere vacuumed to a pressure of 10⁻³ Pa level. At the time of cooling, in the case of a cooling rate of 1° C./sec or lower, temperature control was performed in a vacuum by furnace output control. Further, in the case of the cooling rate of 10° C./sec or more, an Ar gas was introduced and the cooling rate was controlled by adjustment of its flow rate.

Further, the observation of the texture in this period was also performed by the same method described in Example 1. Further, magnetometry was performed in the same manner as that in Example 1 in order to evaluate obtained products, and further the Z value was calculated by the X-ray diffraction.

Further, an alloyed ratio in the second layer and a ratio of the α single phase region were defined and obtained as follows.

Plane distribution of the Fe content and plane distribution of the Al content were measured by using an EPMA (Electron Probe Micro-Analysis) method, with a field of view of an L direction 1 mm×the total thickness in an L cross section. First, as the alloyed ratio in the second layer, areas of a region satisfying Fe≦0.5 mass % and Al≧99.5 mass % before and after the heat treatment were obtained. Then, the alloyed ratio of the second layer was defined as (S₀−S)/S₀×100, where an area when Al was applied and the heat treatment was not performed was set to S₀ and an area in the Fe-based metal sheet on which the whole heat treatment was completed was set to S.

Further, the ratio of the α single phase region was defined as (T/T₀)×100, where an area of a cross section of the Fe-based metal sheet after the heat treatment, observed in the L cross section was set to T₀ and an area of a diffused region of the different metal after the heat treatment was set to T. Incidentally, when the second layer was Al, an area of a region satisfying Al≧0.9 mass % was set to T.

Table 8 shows the base metal sheets and conditions of the decarburization and the heat treatment, and shows the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured during the manufacture (after the decarburization annealing) and after the manufacture (after the diffusion treatment), the Z values of the obtained Fe-based metal sheets, the alloyed ratios of the second layers, and evaluation results of the magnetometry.

TABLE 8 BASE DECARBU- DECARBU- DECARBU- C CONTENT ACCUMULATION TEMPERATURE HOLDING MATERIAL SHEET RIZING RIZATION RIZATION DECARBURIZED AFTER DEGREE OF {200} FERRITE- INCREASING TEMPERATURE C CONTENT THICKNESS ATMOS- TEMPERATURE TIME REGION DECARBURIZATION PLANES AFTER FORMING RATE T1 No. mass% μm PHERE ° C. MINUTE μm mass% DEBARBURIZATION ELEMENT ° C./ sec ° C. 201 0.050 10 STRONG 800 1 9 0.010 26 Al 0.5 1000 202 0.050 100 WEAK+ 800 3 12 0.011 24 Al 0.5 1000 STRONG 203 0.050 250 WEAK+ 800 5 14 0.015 25 Al 0.5 1000 STRONG 204 0.050 500 WEAK+ 800 15 22 0.018 21 Al 0.5 1000 STRONG 205 0.050 750 WEAK+ 800 30 31 0.018 29 Al 0.5 1000 STRONG 206 0.050 100 STRONG 800 0.1 6 0.008 23 Al 0.5 1000 207 0.050 100 WEAK+ 800 250 36 0.017 26 Al 0.5 1000 STRONG 208 0.050 500 WEAK+ 800 550 49 0.017 38 Al 0.5 1000 STRONG 209 0.050 200 WEAK+ 800 10 18 0.008 26 Al 0.1 950 STRONG 210 0.050 200 WEAK+ 800 10 18 0.009 27 Al 1 1000 STRONG 211 0.050 200 WEAK+ 800 10 19 0.008 26 Al 5 1000 STRONG 212 0.050 200 WEAK+ 800 10 18 0.010 25 Al 10 1000 STRONG 213 0.050 200 WEAK+ 800 10 17 0.008 26 Al 20 1000 STRONG 214 0.050 200 WEAK+ 800 10 18 0.009 26 Al 0.5 950 STRONG 215 0.050 200 WEAK+ 800 10 18 0.009 27 Al 0.5 1050 STRONG 216 0.050 200 WEAK+ 800 10 16 0.010 25 Al 0.5 1200 STRONG 217 0.050 150 WEAK+ 800 8 15 0.007 28 Al 0.5 1000 STRONG 218 0.050 150 WEAK+ 800 8 14 0.006 29 Al 0.5 1000 STRONG 219 0.050 150 WEAK+ 800 8 16 0.007 30 Al 0.5 1000 STRONG 220 0.050 150 WEAK+ 800 8 14 0.007 29 Al 0.5 1000 STRONG 221 0.050 150 WEAK+ 800 8 14 0.007 29 Al 0.5 1000 STRONG 222 0.050 150 WEAK+ 800 8 15 0.007 28 Al 0.5 1000 STRONG 223 0.050 150 WEAK+ 800 8 16 0.006 30 Al 0.5 1000 STRONG 224 0.050 300 WEAK+ 800 15 21 0.011 22 Al 0.5 1000 STRONG 225 0.050 300 WEAK+ 800 15 22 0.009 21 Al 0.5 1000 STRONG 226 0.050 300 WEAK+ 800 15 22 0.009 22 Al 0.5 1000 STRONG 227 0.050 8 STRONG 800 1 8 0.010 26 Al 0.5 950 228 0.050 100 STRONG 800 0.05 1 0.050 17 Al 0.5 1000 229 0.050 100 WEAK+ 800 60 59 0.003 28 Al 0.5 1000 STRONG 230 0.050 100 WEAK+ 800 18 23 0.010 26 NONE 0.5 1000 STRONG 231 0.050 100 WEAK+ 800 18 25 0.011 24 Al 0.5 900 STRONG 232 0.050 100 WEAK+ 800 18 26 0.009 27 Al 0.5 1350 STRONG 233 0.050 100 WEAK+ 800 18 25 0.009 23 Al 0.5 1000 STRONG 234 0.050 100 WEAK+ 800 18 24 0.010 25 Al 0.5 1000 STRONG 235 0.050 100 WEAK+ 800 18 26 0.009 28 Al 0.5 1000 STRONG ACCUMULATION ACCUMULATION HOLDING COOLING DEGREE OF {200} DEGREE OF {222} B50/Bs TIME RATE (SO⁻ S) T/ PLANES AFTER PLANES AFTER OF ΔB No. MINUTE ° C./ sec /SO × 100 TO × 100 DIFFUSION DIFFUSION Z PRODUCT T NOTE 201 5 100 79 64 54 16 124 0.892 0.042 INVENTION EXAMPLE 201 202 5 100 65 50 42 28 56 0.864 0.057 INVENTION EXAMPLE 202 203 5 100 52 43 36 24 8.9 0.842 0.098 INVENTION EXAMPLE 203 204 5 100 39 32 36 20 11 0.852 0.091 INVENTION EXAMPLE 204 205 5 100 37 30 37 22 15 0.859 0.085 INVENTION EXAMPLE 205 206 5 100 66 49 46 24 69 0.893 0.042 INVENTION EXAMPLE 206 207 5 100 64 44 41 22 42 0.865 0.054 INVENTION EXAMPLE 207 208 5 100 31 26 32 27 3.1 0.833 0.101 INVENTION EXAMPLE 208 209 1 100 61 50 41 23 39 0.859 0.083 INVENTION EXAMPLE 209 210 1 100 59 48 39 26 25 0.865 0.071 INVENTION EXAMPLE 210 211 5 100 62 51 42 18 68 0.872 0.045 INVENTION EXAMPLE 211 212 5 100 60 46 34 25 4.3 0.851 0.096 INVENTION EXAMPLE 212 213 5 100 58 50 49 11 76 0.897 0.038 INVENTION EXAMPLE 213 214 5 100 59 48 36 23 16 0.845 0.084 INVENTION EXAMPLE 214 215 5 100 60 44 48 18 82 0.896 0.021 INVENTION EXAMPLE 215 216 5 100 57 46 60 9 148 0.904 0.016 INVENTION EXAMPLE 216 217 0.5 100 29 24 34 24 3.5 0.835 0.115 INVENTION EXAMPLE 217 218 10 100 61 47 73 6 175 0.953 0.008 INVENTION EXAMPLE 218 219 30 100 76 55 62 11 152 0.913 0.011 INVENTION EXAMPLE 219 220 60 100 81 68 58 14 135 0.901 0.018 INVENTION EXAMPLE 220 221 120 100 96 75 52 15 112 0.899 0.021 INVENTION EXAMPLE 221 222 550 100 100 74 59 10 139 0.908 0.016 INVENTION EXAMPLE 222 223 4500 100 100 76 55 12 131 0.895 0.018 INVENTION EXAMPLE 223 224 10 0.1 79 64 63 8 162 0.918 0.011 INVENTION EXAMPLE 224 225 10 10 51 43 72 4 189 0.958 0.007 INVENTION EXAMPLE 225 226 10 450 55 37 68 6 158 0.954 0.009 INVENTION EXAMPLE 226 227 1 100 100 100 18 31 0.3 0.789 0.108 COMPARITIVE EXAMPLE 201 228 10 100 65 41 24 37 1.2 0.785 0.110 COMPARITIVE EXAMPLE 202 229 10 100 62 38 37 24 19 0.857 0.086 INVENTION EXAMPLE 227 230 10 100 0 0 19 31 1.4 0.778 0.105 COMPARITIVE EXAMPLE 203 231 10 100 46 37 22 27 0.9 0.768 0.104 COMPARITIVE EXAMPLE 204 232 10 100 78 51 38 20 14 0.832 0.095 INVENTION EXAMPLE 228 233 6000 100 100 76 38 26 14 0.842 0.091 INVENTION EXAMPLE 229 234 10 0.05 85 70 37 25 16 0.845 0.089 INVENTION EXAMPLE 230 235 10 550 34 27 38 26 15 0.841 0.089 INVENTION EXAMPLE 231

As shown in Table 8, in each of present invention examples, it was possible to confirm that the magnetic flux density difference ΔB becomes a small value as compared to comparative examples and a high magnetic flux density can be obtained thoroughly in the in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.80 or more is obtained.

Further, in the present invention examples, it was possible to confirm that the alloyed ratio and the ratio of the α single phase region can be controlled by the combination of the decarburized depth of the base metal sheet, the temperature increasing rate, the holding temperature after the heating, and the holding time, and the Fe-based metal sheet having an excellent magnetic property can be obtained.

Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and the ratio of the α single phase region to the L cross section is 1% or more.

In contrast to this, for example, in the case of the insufficient decarburized region as in a comparative example 201, in the case of using no metal for the second layer as in a comparative example 203, and in the case of not heating to a temperature of the A3 point or higher as in a comparative example 204, it was not possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly as in the present invention examples. Further, even when the temperature was increased to a higher temperature and the holding time was made longer as in present invention examples 228 and 229, the similar effect was able to be obtained, but the significant effect did not appear.

Example 4

In this example, as the ferrite-forming element, Sn, Al, Si, Ti, Ga, Ge, Mo, V, Cr, or As was applied to the second layer, and the relationship between the case where demanganization was performed in addition to decarburization and an accumulation degree of {200} planes was examined.

First, base metal sheets containing six types of components F to K shown in Table 9 below and having a balance being composed of Fe and inevitable impurities were prepared. Ingots were each melted by vacuum melting to then be worked to a predetermined thickness by hot rolling and cold rolling, and the above-described base metal sheets were obtained. Incidentally, the A1 point of each of these base metal sheets was 727° C.

TABLE 9 (MASS%) Ar3 STEEL POINT TYPE °C. C SI Mn Al P N S O F 877 0.03 0.05 0.15 0.0005 0.0001 0.0002 <0.0004 0.0002 G 880 0.03 0.1 0 0.25 0.0004 0.0002 0.0001 <0.0005 0.0001 H 867 0.05 0.05 1.00 0.0003 0.0001 0.0002 <0.0004 0.0002 I 771 0.50 0.1 0.30 0.0004 0.0002 0.0002 <0.0004 0.0002 J 773 0.80 1 .00 0.12 0.0030 0.0020 0.0001 <0.003 0.0001 K 859 0.10 0.30 1.50 0.0030 0.0020 0.0001 <0.003 0.0002

In the hot rolling, the ingots each having a thickness of 230 mm were heated to 1000° C. to be thinned down to a thickness of 50 mm, and hot-rolled sheets were obtained. Then, sheet materials having various thicknesses were cut out from these hot-rolled sheets by machining to then be subjected to the cold rolling, and the base metal sheets each having a thickness falling within a range of 10 μm to 750 μm were manufactured.

At this time, the main phase of each of the obtained base metal sheets at room temperature was an α-Fe phase. Further, as a result of measurement, the A3 point at which the α-γ transformation occurred was temperatures shown in Table 9. Further, by X-ray diffraction, a texture in the α-Fe phase of each of the base metal sheets was measured, and by the previously described method, an accumulation degree of {200} planes and an accumulation degree of {222} planes were obtained. As a result, it was confirmed that at the stage of completion of the cold rolling, of each of the base metal sheets, the accumulation degree of the {200} planes was 19 to 27% and the accumulation degree of the {222} planes was 18 to 25%.

Next, these base metal sheets after this cold rolling each had a material promoting decarburization, or a material promoting decarburization and a material promoting demanganization applied thereto as an annealing separating agent, and were subjected to tight coil annealing or stacked annealing. At this time, the annealing was performed so that depths of the decarburization and the demanganization might become not less than 1 μm nor more than 49 μm. As conditions of the annealing, the temperature was set to 700° C. to 900° C. and the annealing was performed in a reduced pressure atmosphere. Further, a structure and a crystal orientation of a surface layer after completion of the decarburization annealing or the decarburization and demanganization annealing were examined. The measurement of the crystal orientation was performed by the X-ray diffraction method, and the accumulation degree of the {200} planes in the α-Fe phase and the accumulation degree of the {222} planes in the α-Fe phase were obtained.

Next, with respect to each of the base metal sheets after the decarburization annealing or the decarburization and demanganization annealing, both surfaces of each of the base metal sheets were coated with the different metal by using an IP method, a hot dipping method, or a sputtering method to have a thickness of 10 μm in total.

Subsequently, a heat treatment was performed under various conditions by the same method as that used in Example 3, and an experiment was performed in which the state in each of the processes during the manufacture was evaluated. An alloyed ratio of the second layer was defined as (S₀−S)/S₀×100 similarly to Example 3, and assuming that a metal element of the second layer was [M], an area of a region satisfying Fe≦0.5 mass % and [M]≧99.5 mass % was obtained, which was applied to any one of the elements.

On the other hand, a ratio of the α single phase region was also obtained by the same procedure as that in Example 3. However, when the second layer was Sn, T was obtained from an area of a region satisfying Sn≧3.0 mass %, and similarly, in the case of Al, it was obtained from an area of a region satisfying Al≧0.9 mass %. Further, in the case of Si, it was obtained from an area of a region satisfying Si≧1.9 mass %, and in the case of Ti, it was obtained from an area of a region satisfying Ti≧3.0 mass %. Similarly, in the case of Ga, it was obtained from an area of a region satisfying Ga≧4.1 mass %, in the case of Ge, it was obtained from a region satisfying Ge≧6.4 mass %, in the case of Mo, it was obtained from a region satisfying Mo≧3.8 mass %, in the case of V, it was obtained from a region satisfying V≧1.8 mass %, in the case of Cr, it was obtained from a region satisfying Cr≧14.3 mass %, and in the case of As, it was obtained from an area of a region satisfying As≧3.4 mass %.

Table 10 and Table 11 show the base metal sheets and conditions of the heat treatment such as the decarburization, and show the accumulation degrees of the {200} planes and the accumulation degrees of the {222} planes measured during the manufacture (after the decarburization•demanganization annealing) and after the manufacture (after the diffusion treatment), the Z values of the obtained Fe-based metal sheets, the alloyed ratios of the second layers, and evaluation results of the magnetometry.

TABLE 10 DECARBURIZATION AND DECARBURIZED BASE DEMANGANIZATION AND C CONTENT TEMPERATURE HOLDING MATERIAL SHEET ANNEALING DEMANGANIZED AFTER FERRITE- INCREASING TEMPERATURE HOLDING STEEL THICKNESS TEMPERATURE REGION DECARBURIZATION FORMING RATE T1 TIME COOLING RATE No. TYPE μm ° C. μm mass % ELEMENT ° C./sec ° C. MINUTE ° C./sec 236 F 150 800 21 0.008 Sn 0. 5 1000 5 100 237 G 150 800 26 0. 010 Sn 0. 5 1000 5 100 238 H 150 800 23 0.009 Sn 0. 5 1000 5 100 239 I 150 800 24 0. 011 Sn 0. 5 1000 5 100 240 J 150 800 21 0. 009 Sn 0. 5 1000 5 100 241 K 150 800 26 0.009 Sn 0. 5 1000 5 100 242 F 10 800 4 0. 010 Al 0. 5 1000 5 100 243 F 100 800 12 0.011 Al 0. 5 1000 5 100 244 F 250 800 14 0. 015 Al 0. 5 1000 5 100 245 F 500 800 22 0. 018 Al 0. 5 1000 5 100 246 F 750 800 31 0.018 Al 0. 5 1000 5 100 247 G 150 700 10 0. 008 Al 0. 5 1000 5 100 248 G 150 900 24 0. 017 Al 0. 5 1000 5 100 249 H 200 800 49 0.011 Al 0. 5 1000 5 100 250 H 200 800 6 0.014 Al 0. 5 1000 5 100 251 I 100 800 14 0. 006 Al 0. 5 1000 5 100 252 I 100 800 15 0. 014 Si 0. 5 1000 5 100 253 I 100 800 15 0.014 Zn 0. 5 1000 5 100 254 I 200 800 18 0. 008 Ti 0. 5 1000 5 100 255 I 200 800 18 0. 009 Ga 0. 5 1000 5 100 256 I 200 800 19 0. 008 Co 0. 5 1000 5 100 257 I 200 800 18 0. 010 Mb 0. 5 1000 5 100 258 I 200 800 17 0.008 V 0. 5 1000 5 100 259 I 200 800 18 0. 009 Cr 0. 5 1000 5 100 260 I 200 800 18 0. 009 As 0. 5 1000 5 100 261 J 200 800 16 0. 010 Al 0. 1 1000 5 100 262 J 150 800 15 0.007 Al 1 1000 5 100 ACCUMULATION ACCUMULATION ACCUMULATION ACCUMULATION DEGREE OF DEGREE OF DEGREE OF DEGREE OF {200} PLANES {222} PLANES {200} PLANES {222} PLANES AFTER AFTER AFTER AFTER (SO − S) T/ ΔB No. ANNEALING ANNEALING DIFFUSION DIFFUSION SO × 100 TO × 100 Z PRODUCT T NOTE 236 36 21 61 14 66 38 126 0.921 0.021 INVENTION EXAMPLE 232 237 34 25 68 12 64 33 135 0.942 0.034 INVENTION EXAMPLE 233 238 33 21 65 13 59 36 129 0.937 0.037 INVENTION EXAMPLE 234 239 35 23 71 8 61 37 187 0.963 0.007 INVENTION EXAMPLE 235 240 37 21 59 17 65 34 113 0.921 0.048 INVENTION EXAMPLE 236 241 34 18 60 12 63 35 123 0.917 0.042 INVENTION EXAMPLE 237 242 35 21 47 17 79 64 64 0.884 0.069 INVENTION EXAMPLE 238 243 34 24 49 13 65 50 56 0.879 0.072 INVENTION EXAMPLE 239 244 35 18 51 11 52 43 63 0.892 0.062 INVENTION EXAMPLE 240 245 31 26 44 19 39 32 52 0.876 0.075 INVENTION EXAMPLE 241 246 39 22 40 18 37 30 23 0.857 0.086 INVENTION EXAMPLE 242 247 24 26 46 14 66 49 35 0.872 0.082 INVENTION EXAMPLE 243 248 36 18 68 9 64 44 142 0.924 0.027 INVENTION EXAMPLE 244 249 38 27 72 4 31 26 167 0.947 0.011 INVENTION EXAMPLE 245 250 23 20 39 15 68 57 5.4 0.853 0.101 INVENTION EXAMPLE 246 251 29 16 40 17 69 59 28 0.857 0.087 INVENTION EXAMPLE247 252 34 24 46 16 66 54 37 0.879 0.083 INVENTION EXAMPLE 248 253 33 26 51 13 67 56 67 0.896 0.067 INVENTION EXAMPLE 249 254 36 22 49 19 61 50 29 0.887 0.076 INVENTION EXAMPLE 250 255 37 24 56 8 59 48 121 0.897 0.031 INVENTION EXAMPLE 251 256 36 23 55 8 62 51 116 0.916 0.036 INVENTION EXAMPLE 252 257 35 26 61 6 60 46 129 0.916 0.026 INVENTION EXAMPLE 253 258 36 19 48 17 58 50 59 0.897 0.068 INVENTION EXAMPLE 254 259 36 28 60 9 59 48 129 0.906 0.039 INVENTION EXAMPLE 255 260 37 15 53 11 60 44 119 0.899 0.042 INVENTION EXAMPLE 256 261 35 22 57 10 57 46 125 0.914 0.029 INVENTION EXAMPLE 257 262 38 17 61 9 29 24 131 0.916 0.013 INVENTION EXAMPLE 258

TABLE 11 DECARBURIZATION BASE AND C CONTENT TEMPERATURE HOLDING MATERIAL SHEET ANNEALING DEMANGANIZATION AFTER FERRITE- INCREASING TEMPERATURE HOLDING COOLING STEEL THINCKNESS TEMPERATURE REGION DECARBURIZATION FORMING RATE T1 TIME RATE No. TYPE μm ° C. μm mass % ELEMENT ° C./sec ° C. MINUTE ° C./sec 263 J 150 800 14 0.006 Al 5 1000 5 100 264 J 150 800 16 0.007 Al 10 1000 5 100 265 J 150 800 14 0.007 Al 20 1000 5 100 266 K 150 800 14 0.007 Al 0.5 950 5 100 267 K 150 800 15 0.007 Al 0.5 1250 5 100 268 K 150 800 16 0.006 Al 0.5 1000 0. 5 100 269 K 300 800 21 0.011 Al 0.5 1000 10 100 270 K 300 800 22 0.009 Al 0.5 1000 30 100 271 K 300 800 22 0.009 Al 0.5 1000 60 100 272 K 300 800 21 0.008 Al 0.5 1000 120 100 273 K 300 800 23 0.009 Al 0.5 1000 550 100 274 K 300 800 21 0.008 Al 0.5 1000 4500 100 275 G 300 800 21 0.008 Al 0.5 1000 5 0. 1 276 G 300 800 21 0.008 Al 1.5 1000 5 10 277 G 300 800 21 0.008 Al 2.5 1000 5 450 278 F 8 800 8 0.010 Al 0.5 950 1 100 279 G 100 650 21 0.050 Al 0.5 1000 10 100 280 G 100 950 41 0.003 Al 0.5 1000 10 100 281 H 100 800 1 0.010 Al 0.5 1000 10 100 282 H 200 900 69 0.011 Al 0.5 1000 10 100 283 K 100 800 26 0.009 NONE 0.5 1000 10 100 284 I 100 800 25 0.009 Al 0.5 765 10 100 285 I 100 800 24 0.010 Al 0.5 1350 10 100 286 J 100 800 26 0.009 Al 0.5 1000 6050 100 287 J 100 800 24 0.008 Al 0.5 1000 10 0.05 288 J 100 800 26 0.010 Al 0.5 1000 10 500 ACCUMULATION ACCUMULATION ACCUMULATION ACCUMULATION DEGREE OF DEGREE OF DEGREE OF DEGREE OF {200} PLANES {222} PLANES {200} PLANES {222} PLANES AFTER AFTER AFTER AFTER (SO − S) T/ ΔB No. ANNEALING ANNEALING DIFFUSION DIFFUSION SO × 100 TO × 100 Z PRODUCT T NOTE 263 39 23 75 7 61 47 189 0.975 0.006 INVENTION EXAMPLE 259 264 30 19 42 16 76 55 43 0.864 0.064 INVENTION EXAMPLE 260 265 29 14 38 16 81 68 16 0.846 0.098 INVENTION EXAMPLE 261 266 29 24 40 11 96 75 21 0.853 0.092 INVENTION EXAMPLE 262 267 28 23 36 17 100 74 8.3 0.839 0.103 INVENTION EXAMPLE 263 268 30 25 67 7 100 76 164 0.943 0.010 INVENTION EXAMPLE 264 269 22 29 43 15 79 64 53 0.872 0.059 INVENTION EXAMPLE 265 270 21 30 41 18 51 43 43 0.867 0.063 INVENTION EXAMPLE 266 271 22 28 38 16 55 37 12 0. 843 0.096 INVENTION EXAMPLE 267 272 28 21 64 5 56 34 158 0.929 0.009 INVENTION EXAMPLE 268 273 27 19 73 3 53 38 168 0. 968 0.007 INVENTION EXAMPLE 269 274 22 31 79 7 55 42 198 0.978 0.005 INVENTION EXAMPLE 270 275 29 25 51 4 54 38 123 0.895 0.036 INVENTION EXAMPLE 271 276 26 24 52 6 51 44 128 0.896 0.034 INVENTION EXAMPLE 272 277 24 25 45 12 53 31 73 0.879 0.053 INVENTION EXAMPLE 273 278 26 23 47 19 100 100 86 0.876 0.049 INVENTION EXAMPLE 274 279 17 19 21 28 65 41 1.3 0.778 0.123 COMPARITIVE EXAMPLE 205 280 14 14 19 22 62 38 0.8 0.779 0.113 COMPARITIVE EXAMPLE 206 281 11 24 23 14 66 31 0.9 0.782 0.109 COMPARITIVE EXAMPLE 207 282 24 23 57 18 61 37 135 0.905 0.037 INVENTION EXAMPLE 275 283 27 25 12 11 0 0 0.6 0.765 0.096 COMPARITIVE EXAMPLE 208 284 23 29 25 31 57 21 1.2 0.786 0.109 COMPARITIVE EXAMPLE 209 285 25 16 48 22 85 70 68 0.875 0.052 INVENTION EXAMPLE 276 286 28 24 46 15 92 73 63 0.881 0.054 INVENTION EXAMPLE 277 287 22 26 38 17 84 69 12 0.852 0.089 INVENTION EXAMPLE 278 288 26 22 39 16 63 31 21 0.859 0.086 INVENTION EXAMPLE 279

In each of present invention examples, it was possible to confirm that the magnetic flux density difference ΔB becomes a small value as compared to comparative examples and a high magnetic flux density is obtained thoroughly in the in-plane circumferential direction. Further, in these Fe-based metal sheets, it was possible to confirm that an excellent magnetic property in which the value of B50/Bs is 0.80 or more is obtained.

Further, in the present invention examples, as shown in Table 10 and Table 11, it was possible to confirm that the {200} plane in the α-Fe phase is likely to be highly accumulated at each of the stages of the heat treatment.

Further, an L cross section of each of the present invention examples was observed, and thereby it was confirmed that the α single phase region made of the α single phase based component exists in at least a partial region including the surfaces and the ratio of the α single phase region to the L cross section is 1% or more.

In contrast to this, for example, in the case of the insufficient decarburized and demanganized region as in a comparative example 207, in the case of using no metal for the second layer as in a comparative example 208, and in the case of not heating to a temperature of the A3 point or higher as in a comparative example 209, it was not possible to obtain a high magnetic flux density in the in-plane circumferential direction thoroughly as in the present invention examples, and consequently, an obtained magnetic property was also poor. Even when the temperature was increased to a higher temperature and the holding time was made longer as in present invention examples 276 and 277, the similar effect was able to be obtained, but the significant effect did not appear.

In the foregoing, the preferred embodiments of the present invention have been described in detail, but the present invention is not limited to such examples. It is apparent that a person having common knowledge in the technical field to which the present invention belongs is able to devise various variation or modification examples within the range of technical ideas of the present invention, and it should be understood that they also belong to the technical scope of the present invention as a matter of course.

INDUSTRIAL APPLICABILITY

The Fe-based metal sheet of the present invention is suitable for magnetic cores and the like of transformers and the like using a silicon steel sheet, and can contribute to downsizing of these magnetic cores and reduction in energy loss. 

1. An Fe-based metal sheet, comprising: less than 0.2 mass % C and having a composition that is capable of causing an α-γ transformation, wherein: a ferrite-forming element being Si is alloyed on a partial or whole region of the Fe-based metal sheet, and when intensity ratios of respective {001}<470>, {116}<6 12 1>, and {223}<692> directions in a sheet plane by X-ray diffraction are set to A, B, and C respectively and Z=(A+0.97B)/0.98C is satisfied, a Z value is not less than 2.0 nor more than 200; wherein at least a partial region including surfaces of the Fe-based metal sheet is an α single phase region that is alloyed with said ferrite-forming element, and a ratio of the α single phase region to a cross section of the Fe-based metal sheet is 1% or more.
 2. The Fe-based metal sheet according to claim 1, wherein: a layer containing said ferrite-forming element is formed on at least one side of surfaces of the Fe-based metal sheet, and a ferrite-forming element that has diffused from part of the layer is alloyed with Fe.
 3. The Fe-based metal sheet according to claim 2, wherein: a thickness of the layer containing said ferrite-forming element is not less than 0.01 μm nor more than 500 μm.
 4. The Fe-based metal sheet according to claim 1, wherein: an accumulation degree of {200} planes is not less than 30% nor more than 99%, and an accumulation degree of {222} planes is not less than 0.01% nor more than 30%, and said accumulation degree of {200} planes is represented by Expression (1) below, and said accumulation degree of {222} planes is represented by Expression (2) below: accumulation degree of {200} planes=[{i(200)/I(200)}/Σ{i(hkl)/I(hkl)}]×100   Expression (1) accumulation degree of {222} planes=[{i(222)/I(222)}/Σi(hkl)/I(hkl)}]×100   Expression (2) wherein i (hkl) is an actually measured integrated intensity of {hkl} planes in a surface of the Fe-based metal sheet, and I (hkl) is a theoretical integrated intensity of the {hkl} planes in a sample having a random orientation, and 11 kinds of planes of {110}, {200}, {211}, {310}, {222}, {321}, {411}, {420}, {332}, {521}, and {442} are used as the {hkl} planes.
 5. The Fe-based metal sheet according to claim 1, wherein: a thickness of the Fe-based metal sheet is not less than 10 μm nor more than 6 mm.
 6. The Fe-based metal sheet according to claim 1, wherein: the α single phase region is formed on a front surface side and a rear surface side of the Fe-based metal sheet, and a crystal grain straddling the α single phase region on the front surface side and the α single phase region on the rear surface side is formed. 